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The role of nanotwins and grain boundary plane in the thermal, corrosion, and sensitization behavior of nanometals

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The role of nanotwins and grain boundary plane in the thermal, corrosion, and sensitization behavior of nanometals

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Content THE ROLE OF NANOTWINS AND GRAIN BOUNDARY PLANE IN THE THERMAL, CORROSION, AND SENSITIZATION BEHAVIOR OF NANOMETALS by Yifu Zhao A Dissertation Presented to the FACULTY OF THE GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA In Partial Fulfillment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (MECHANICAL ENGINEERING) August 2014 i To my parents, Fuchen Zhao and Jianping Wang, whom I deeply love ii Acknowledgements I would like to thank my advisors Prof. Andrea Hodge and Prof. Michael Kassner for their abundant guidance throughout the five years of my PhD study. Specifically, I want to thank Prof. Hodge for providing the research directions, equipment, and kind exhortations, so that my experiments could be completed in a timely fashion. Many thanks must also be given to Prof. Kassner who picked me into the PhD program and patiently trained me to become an earnest researcher. I also like to thank Prof. Veronica Eliasson, Prof. Edward Goo, Prof. John Platt, and Prof. Pin Wang who generated many insightful discussions and advices during my qualifying exam and PhD defense. I am very grateful to my labmates. Particularly, thank you to Mikhail Polyakov for the genuine support and help on the FIB lift-out and TEM characterization experiments; thank you to Timothy Furnish and Leonardo Estrada for kindly providing the sputtered Cu and Al samples; thank you to I-Chung Cheng for the patient support and suggestions throughout my experiments. My research would not have been that efficient without the many discussions, advices, and help from my labmates. Moreover, I would also like to thank Yuzheng Zhang and Jittraporn Wongsa- Ngam for the insightful ideas and consistent support given to me. Many thanks should also be given to John Curulli and Matthew Mecklenburg who trained and helped me to use the characterization equipment in the CEMMA of USC. iii Finally, warm thanks must be given to my father Fuchen Zhao and my mother Jianping Wang, who always support me, trust me, and love me. I also want to thank Jesus Christ for giving me the spirit, instructing me to do honest work, and offering me the confidence and wisdom to peacefully carry out the research. iv Table of Contents Acknowledgements ........................................................................................................ ii Abstract ........................................................................................................................ vii Chapter 1 Introduction ...................................................................................................1 Chapter 2 Background ...................................................................................................5 2.1 Grains and grain boundaries ..................................................................................5 2.2 Grain boundary engineering ................................................................................ 12 2.3 Nanotwinned metals and their mechanical properties ........................................... 15 2.3.1 Nanotwinned metals ....................................................................................... 15 2.3.2 Mechanical properties of nanotwinned metals................................................ 20 2.4 Thermal stability ................................................................................................. 22 2.4.1 The driving force for grain growth ................................................................. 23 2.4.2 Grain growth in nanotwinned metals ............................................................. 24 2.5 Corrosion behavior .............................................................................................. 26 2.5.1 The role of grain boundary in corrosion resistance ........................................ 27 2.5.2 Corrosion behavior of nanotwinned metals .................................................... 30 2.6 Sensitization of Al-Mg alloys ............................................................................... 32 2.6.1 Sensitization and its consequences ................................................................. 32 2.6.2 Model of β phase growth ................................................................................ 36 2.6.3 Grain boundary effects................................................................................... 38 Chapter 3 Characterization and test methods ................................................................ 41 3.1 Characterization methods .................................................................................... 41 3.1.1 X-ray diffraction (XRD) ................................................................................. 41 3.1.2 Scanning electron microscope (SEM) ............................................................. 43 3.1.3 Electron backscatter diffraction (EBSD) ........................................................ 46 3.1.4 Focused ion beam (FIB) ................................................................................ 50 3.1.5 Transmission electron microscope (TEM) ...................................................... 54 3.2 Test methods ....................................................................................................... 57 v 3.2.1 Heat-treatments ............................................................................................. 57 3.2.2 Polarization ................................................................................................... 58 3.2.3 Immersion corrosion ...................................................................................... 64 3.2.4 Chemical etching ........................................................................................... 64 Chapter 4 The effect of nanotwins on the thermal stability of Cu ................................. 68 4.1 Experimental procedures ..................................................................................... 68 4.2 As-prepared samples ........................................................................................... 70 4.3 Grain growth at elevated temperature .................................................................. 71 4.4 Grain boundary migration.................................................................................... 76 4.5 Grain orientation ................................................................................................. 79 4.6 Conclusions ......................................................................................................... 84 Chapter 5 The effect of nanotwins on the corrosion behavior of Cu ............................. 86 5.1 Experimental procedures ..................................................................................... 86 5.2 As-prepared samples ........................................................................................... 89 5.3 Polarization .......................................................................................................... 92 5.3.1 Polarization curve analysis ............................................................................ 92 5.3.2 Passive layer morphology .............................................................................. 98 5.3.3 Corrosion paths ........................................................................................... 100 5.4 Immersion corrosion.......................................................................................... 101 5.5 Composition of passive layers ........................................................................... 104 5.6 Conclusions ....................................................................................................... 106 Chapter 6 The effect of grain boundary plane orientation on the sensitization of an Al- Mg alloy ...................................................................................................................... 108 6.1 Experimental procedures ................................................................................... 108 6.2 As-prepared samples ......................................................................................... 110 6.3 Sensitization behavior of as-received (AR) 5456 ............................................... 112 6.4 Sensitization behavior of as-sputtered (AS) 5456............................................... 116 6.5 Conclusions ....................................................................................................... 119 Chapter 7 Summary and future outlook ...................................................................... 120 vi Appendix A Abbreviations ........................................................................................ 124 Appendix B Further analysis of the anodic Tafel region ............................................. 126 Appendix C Measurement of grain boundary plane orientations ................................ 128 Appendix D List of figures and tables ........................................................................ 136 References ................................................................................................................... 146 vii Abstract Grain boundaries play a significant role in the properties of polycrystalline metals, and the objective of grain boundary engineering is to incorporate special grain boundaries into materials to achieve better properties. While a special grain boundary is conventionally defined as a boundary with coincidence site lattice  ≤ 29, it should be noted that the grain boundary plane can also largely affect grain boundary structure and properties. Nanotwins in particular, which are bundles of coherent 3 boundaries, have extremely low energy and improved properties compared with ordinary grain boundaries. While it is well-studied that nanotwins could enhance the mechanical properties of a metal, the effect of nanotwins on the thermal stability and corrosion resistance still needs to be explored and is the subject of this research. In addition to nanotwins which have {111} oriented boundary planes, this study also explored how differently oriented grain boundary planes could affect the sensitization behavior of an Al-Mg alloy. For the thermal stability study, sputtered Cu foils with two types of microstructures, highly nanotwinned and non-nanotwinned columnar grains, were sequentially heat-treated at 200, 300, and 400 C. The highly nanotwinned Cu (grain size,  700 nm) remained thermally stable up to 300 C, whereas the non-nanotwinned Cu (grain size,  400 nm) had a rapid grain growth even at 200 C. Additionally, the effect of nanotwins on corrosion behavior was evaluated through testing Cu samples containing various fractions of nanotwins: a) highly nanotwinned columnar grains, b) partially nanotwinned columnar grains, c) non-nanotwinned columnar grains, and d) viii microcrystalline grains. These samples were tested in 3.5% NaCl solution (pH ~ 8.0) by linear polarization, potentiodynamic polarization, and immersion corrosion methods. It was found that highly nanotwinned Cu had a lower corrosion current density and a more protective passive layer compared with non-nanotwinned, partially nanotwinned, and microcrystalline Cu samples. The improved thermal stability and corrosion resistance of highly nanotwinned Cu were attributed to the ordered structure of nanotwins, the special grain boundary network, and the {111} texture due to the presence of nanotwins. While a {111} oriented coherent boundary plane of nanotwins can enhance the thermal and corrosion behavior, this study further identified grain boundary plane orientations and explored their roles in the sensitization behavior of conventional and sputtered Al 5456 samples. Both types of samples were heated at 175 C and subsequently etched in 10% H 3 PO 4 in order to reveal the Mg-rich β phase. The orientations of selected grain boundary planes were measured and related to their corresponding β phase thicknesses. It was found that grain boundaries with plane orientations close to {110} appear particularly vulnerable to β precipitation. Moreover, the sputtered Al 5456 had much higher resistance to β precipitation than the conventional Al 5456, which could be due to its special columnar grain boundary plane orientations. This study highlights the potential of incorporating nanotwins and altering grain boundary plane orientations in order to improve thermal stability, corrosion resistance, and sensitization behavior. 1 Chapter 1 Introduction Thermal stability and corrosion resistance are two important material properties for industrial applications. At elevated temperatures, metals with inadequate thermal stability could experience grain growth which could lead to property changes such as a decrease in strength. In a corrosive environment, metals are subject to surface dissolution and deterioration, leading to degradation in mechanical properties. Particularly, for Al- Mg (5xxx series) alloys with more than 3 wt% Mg, the Mg atoms could diffuse toward grain boundaries (GBs) and form intergranular β phase (Al 3 Mg 2 ). Since the intergranular β phase is anodic to the Al-Mg matrix, it will be preferentially corroded, resulting in intergranular corrosion and stress corrosion cracking [1, 2]. This phenomenon is called the “sensitization” of Al-Mg alloys. Given that the thermal, corrosion, and sensitization behavior may lead to the deterioration and failure of materials used especially in elevated-temperature or corrosive environments, further studies and improvements are needed. While there are various approaches to alter these behaviors, this thesis focuses on grain boundary engineering (GBE) and explores ways to use GBs to improve thermal, corrosion, and sensitization properties. The objective of GBE is to incorporate special GBs into the materials in order to improve the overall properties. Although special GBs are conventionally defined as GBs with coincidence site lattice (CSL)  ≤ 29 [3], it has been reported [4-6] that the GB plane orientation also plays a significant role. Nevertheless, coherent twin boundaries (CTBs), which are 3 GBs with the GB planes coherently oriented in {111}, are 2 considered to be special GBs [7]. In this study, magnetron sputtering was used to synthesize nanotwinned (nt) Cu samples which contain CTBs with nanoscale spacings. Due to the presence of high-density CTBs (or nt GBs), these nt Cu samples are a special kind of GBE material. The effect of the nanotwins on the thermal stability and corrosion resistance was studied. Additionally, this study further explored the sensitization behavior of conventional and sputtered Al-Mg alloys in order to identify the effect of GB plane orientation on β precipitation. A study by Anderoglu et al. [8] on nt Cu (twin thickness  4 nm, grain size  40 nm) at 800 C reported modest nt growth and an order of magnitude grain growth. However, as pointed out by Zhang and Misra [9], the superior thermal stability of nt materials requires further study and must include additional factors such as texture. The uniqueness of this thesis is that the thermal stability of highly nt Cu was evaluated in the perspective of GB character and texture. Specifically since the GB character affects GB migration, this aspect becomes more critical with the presence of nanotwins. Some studies [8-12] have shown that the presence of low-energy CTBs enhanced the thermal stability, so it is expected that large amounts of nt GBs would lead to better overall thermal stability. In this work, the thermal stability of highly nt Cu with ultra-fine grain (ufg) and twin thickness of  40 nm is compared with ufg Cu without nanotwins. The samples were heat-treated at regular intervals until the nt structure was no longer stable, and GB migration and texture changes were characterized before and after heat- treatments. 3 There have been no published corrosion studies in nt metals with low stacking fault energy (SFE, ϒ SFE ) such as Ag (ϒ SFE = 16 – 22 mJ/m 2 ) and Cu (ϒ SFE = 35 – 78 mJ/m 2 ) [13, 14]. However, there are some studies [15-18] on the corrosion behavior of sputtered Al and electrodeposited Ni samples containing nanotwins, although these types of materials tend not to twin due to the high SFE of Al (ϒ SFE = 135 – 200 mJ/m 2 ) and Ni (ϒ SFE = 125 – 300 mJ/m 2 ) [13, 14]. It is therefore difficult to assess whether the samples in the Ni and Al studies were highly nanotwinned, especially since the manuscripts are mainly focused on the electrochemical aspect of corrosion rather than the microstructural characterization. However, these studies did provide an initial assessment about the possible effects of nanotwins on corrosion. For example, Meng et al. [15-17] studied the electrochemical behavior of Ni with nanotwins in borate buffer solution with and without NaCl addition, and found that its passive layer was thinner, less defective, and had a higher pitting corrosion resistance compared with both cast Ni and industrial electrodeposited (presumably without nanotwins) Ni. It was also reported [18] that Al with nanotwins had a lower corrosion current density and a higher pitting resistance than cast Al in acidic NaCl solution. The corrosion behavior of nt Cu samples was studied by both polarization and immersion corrosion tests in artificial seawater (3.5% NaCl, pH  8.0) in this work. The GB character [19], grain size [20], and grain orientation [21] could strongly affect the corrosion properties of a metal. Therefore, it is crucial to sufficiently characterize the microstructures with and without nanotwins in order to explore the effect of nanotwins on the corrosion properties. Three types of Cu foils having different fractions of nt columnar grains (or columnar grains with nanotwins) were studied and 4 compared to microcrystalline (mc) Cu. The microstructure was characterized before and after corrosion. The passive layer compositions were also identified in order to relate to the observed passivation behavior. GB plane orientations, in addition to the nt GB which has a {111} GB plane orientation, were identified and their effects on the sensitization behavior of Al 5456 was explored. From the perspective of GB misorientation, Davenport et al. [22] reported that low-angle GBs are resistant to sensitization, while Kaigorodova [23] and Scotto D’Antuono et al. [24] observed β precipitation along low-angle GBs. Nevertheless, these previous studies indicated that GB plane orientation could play a significant role in β precipitation [22-24]. There is no study that directly characterized the GB plane orientations and related them to the extent of β precipitation. Therefore, this study focuses on the role of GB plane orientation in the β precipitation along both low-angle and high-angle GBs. 5 Chapter 2 Background 2.1 Grains and grain boundaries Most metals are polycrystalline, which are comprised of grains. Each grain is formed by periodic repetition of a single crystal structure. The common crystal structures in metals are body centered cubic (bcc), face centered cubic (fcc), and hexagonal close packed (hcp) as shown in Fig.1. For example: Ta, W have bcc structure; Cu, Al, Ni have fcc structure; Mg, Ti, Zn have hcp structure. The research presented in this study will mainly focus on fcc structured Cu and Al. Figure 1 Atom positions of three crystal structures: bcc, fcc, hcp [25]. Grains in one metal commonly have a single crystal structure, but different crystal orientations. Fig.2 shows the three dimensional grain distribution of a Ni-based superalloy block studied by electron backscatter diffraction (EBSD, see Section 3.1.3 for more details) [26]. Each color, which represents a specific crystal orientation, constitutes a grain. The interface separating these grains is a GB. 6 Figure 2 3D reconstruction of Ni-based superalloy by EBSD [26]. A GB can be characterized by five variables. Three variables are to define GB misorientation which includes misorientation angle (θ) and misorientation axis, and the other two variables characterize GB plane orientation. Based on the misorientation angle, GBs can be categorized into low-angle GB and high-angle GB. Low-angle GB has a misorientation angle smaller than 10° - 15°, and is consisted of a periodic arrangement of dislocations [27]. However, for high-angle GB (θ > 10° - 15°), its dislocation cores are overlapped, so the structure of a high-angle GB is a lot complex than that of a low-angle GB. The soap bubble model in Fig.3 illustrates that low-angle GB has an array of dislocations separated from each other, while high-angle GB has overlapped dislocations. 7 Figure 3 The soap bubble model of low-angle GB and high-angle GB [27]. Furthermore, another way to categorize GBs is by CSL. All the low-angle GBs and part of the high-angle GBs can be categorized into CSL GBs. The uniqueness of CSL GB is that some atom positions in the CSL GB coincide with the ideal atom positions of the neighboring grains. This coincidence of atom positions can lead to a more ordered GB structure and thus better properties in corrosion, creep, and fatigue [3]. Coincidence sites are the atom positions shared by both neighboring grains, and they form CSL. Since the density of coincident sites can reflect GB structure, a parameter   inverse density of coincident sites  is used to classify different CSL GBs as presented in Equation 1 [27]: Equation 1 8 Table 1 Misorientation angles and axes of special GBs in cubic-crystal metals [28].  Type of special GBs Misorientation angle (°) Misorientation axis 1 0 − 3 60 <111> 5 36.86 <100> 7 38.21 <111> 9 38.94 <110> 11 50.47 <110> 13a 22.62 <100> 13b 27.79 <111> 15 48.19 <210> 17a 28.07 <100> 17b 61.92 <221> 19a 26.53 <110> 19b 46.83 <111> 21a 21.78 <111> 21b 44.41 <211> 23 40.45 <311> 25a 16.26 <100> 25b 51.68 <331> 27a 31.59 <110> 27b 35.43 <210> 29a 43.60 <100> 29b 46.40 <221> Lower value indicates a higher density of coincident sites, a higher lattice matching between the neighboring grains, and thus a more ordered GB structure. Since low- GBs ( ≤ ) are reported [3] to have special mechanical and chemical properties, they are conventionally known as “special GBs”; on the other hand, those disordered high-energy GBs are known as “ordinary GBs”. Table 1 lists the misorientation angles and axes of special GBs for cubic crystal structured metals. Specifically, 1 GBs 9 represent low-angle GBs, and 3 are twin boundaries which can be annealing twins, deformation twins, or growth twins depending on the formation process. The atomic arrangement of a5 GB is shown in Fig.4: The red and yellow dots illustrate the atomic arrangements of two neighboring grains, and the green dots are the coincidence sites where the atoms of neighboring grains overlap. For each coincidence site lattice (the green dots), there are five crystal atoms (either red or yellow dots) from each neighboring grain. Therefore this GB is a 5 GB. Figure 4 Schematic graph of a 5 CSL: red and yellow dots represent the atoms of two crystal lattices, and green dots represent the coincident sites [28]. Nevertheless, a definition of special GBs that solely based on values is debatable, as it only considers GB misorientation but ignores the influence of GB plane orientation. Fig.5 shows the computed GB energies with respect to various  values. It 10 can be generalized from this plot that low-  GBs (especially 5) do not necessarily have low energy and ordered GB structure. In addition, even for a given  value, there is a wide variation of GB energies. For example, the coherent 3 GB (or CTB, marked by the black legend in Fig.5), which has a {111} oriented GB plane, has an extremely low GB energy compared with the other 3 GBs. Figure 5 The computed GB energies of Ni for ranging from 3 to 400.  is the inverse density of coincident sites. Different symbols in the inset represent different types of GBs [7]. The reason for the much lower energy of a CTB is its coherent GB plane. Fig.6 shows the atomic arrangements on the two sides of a CTB in a fcc metal, and the coherent GB plane is marked by the vertical line at the center. It can be seen from Fig.6 (a) that there is a mirror relationship between the two sides of the CTB. Specifically, the sequence of {111} planes in fcc metals changes from “ABCABC” to “CBACBA” across the GB plane. Fig.6 (b) illustrates that although it is a fcc metal, there is a hcp structure 11 in the close vicinity of a CTB because of the change in the sequence of {111} planes. Due to this ordered GB structure, CTB has an extremely low GB energy. Figure 6 Schematic atomic arrangement on the two sides of a CTB in fcc metals: (a) shows the {111} planes marked as “A”, “B”, “C” (vertical direction), adapted from [29]; (b) presents the hcp structure of the CTB, adapted from [30]. Therefore, in addition to GB misorientation, GB plane orientation also plays an important role in GB energy and structure. Randle et al. [4-6] further reported that GB plane orientation could significantly affect GB properties, so a GB analysis that considers both GB misorientation and GB plane orientation would give a more comprehensive view of GB structure and its effect on material properties. 12 2.2 Grain boundary engineering The objective of GBE is to incorporate a large amount of special GBs into the microstructure in order to improve material properties. From Section 2.1, we know that special GBs typically have ordered structure and improved properties. In addition, the incorporation of special GBs can interrupt the connections between ordinary GBs and thus enhance corrosion resistance, thermal stability, and overall mechanical properties [3, 6, 31, 32]. Figure 7 GB distributions of the (a) starting base stainless steel, and (b) GBE stainless steel. Black and gray lines represent ordinary and special GBs, respectively [33]. To produce GBE metals, there are generally two approaches: top-down, and bottom-up. Top-down approach is the traditional method that increases the fraction of special GB by processing bulk metals. Thermal mechanical processing is the most popular among the processing methods of top-down approach. Fig.7 (a) and (b) present the GB distributions of stainless steel before and after thermal mechanical processing. 13 We observe from Fig.7 (b) that the fraction of special GBs (gray lines) greatly increases after thermal mechanical processing, and these newly formed special GBs break up the connectivity of ordinary GBs which means that ordinary GBs are more isolated from each other. In addition to the traditional top-down approach, a high fraction of special GBs can also be achieved in metals that are produced atom by atom, which is the bottom-up approach. Specifically, the highly nt Cu samples of this study were produced by magnetron sputtering. Sputtered nt Cu contains large amounts of CTBs which are special GBs defined in terms of all five degrees of GB freedom. Highly nt Cu will be discussed in Section 2.3. Figure 8 Cross-sectional optical micrographs of (a) conventional and (b) GBE alloy 800 (Fe- 35Ni-25Cr, heat-treated at 600°C for 1 hour) after 120-hours exposure to solution containing sulfuric acid and ferric sulfate (ASTM G28 standard) [3]. Because of the ordered structure of special GBs, GBE metals exhibit improved corrosion properties. For example, Fig.8 (a) and (b) present the cross-sectional 14 micrographs of both ordinary and GBE alloy 800 samples after 120-hours exposure to a sulfuric acid and ferric sulfate solution [3]. One can see that while the ordinary alloy 800 experienced severe intergranular corrosion, GBE alloy 800 shows much better intergranular corrosion resistance. Figure 9 Cross-sectional models and optical micrographs showing the propagation of intergranular corrosion through the GB network. In (a) and (b), “R” refers to random or ordinary GBs; “c” and “i” refer to coherent and incoherent 3 GBs, respectively. The thick black lines are corroded GBs, and the thin lines are non-corroded GBs as shown by the legend at bottom right. (a) and (b) are from ref. [19]; (c) is from ref. [34]. The mechanism that special GBs could help to improve the corrosion property is illustrated in Fig.9. It can be observed from Fig.9 (a) and (b) that corrosion can proceed through 9, 27, and ordinary GBs, but they are arrested by two coherent 3 boundaries (or CTBs). The model in Fig.9 (c) also illustrates that the presence of twin ( 3) boundaries could produce low-energy segments in the high-energy GB network, and 15 these low-angle segments could arrest the intergranular corrosion. Therefore, the special GBs of GBE metals interrupt the network of ordinary GBs, and can improve intergranular corrosion resistance. Besides corrosion resistance, it is also reported [3, 32] that GBE metals possess improved thermal stability, creep resistance, and fatigue properties. Therefore, GBE is a promising field to improve industrial metals without much additional cost for the processing. 2.3 Nanotwinned metals and their mechanical properties Nt metals are a newly-emerged group of metals that received much attention due to the presence of nanotwins and the subsequent special properties. This section will introduce the microstructure and production methods of nt metals, and then discuss their mechanical properties. 2.3.1 Nanotwinned metals Nt metal is a special type of metal that contains high-density CTBs within its grains. Since the spacing between the CTBs in nt metals is of nanoscale (typically 1 – 100 nm), it is called nanotwin. Fig.10 illustrates the high-resolution transmission electron microscope (TEM) graph of nanotwins. It can be seen that the twinning plane (or CTB plane) is (1 1 ¯ 1), and the atomic alignments on the two sides of the twinning plane have a mirror relationship. Fig.10 also reveals the nanoscale spacings between 16 CTBs. The purpose of producing nt metals is to improve material properties by these high-density CTBs, so nt metal is a special type of GBE metal. Figure 10 Atomic arrangement of nanotwins studied by high-resolution TEM. The solid white lines indicate (1 1 ¯ 1), (1 ¯ 1 ¯ 1), and (2 0 0) planes, respectively [35]. Since nanotwins in fcc metals are basically the changes in the stacking sequence of the {111} planes as discussed in Section 2.1, they can be synthesized more easily in metals with low SFEs (such as Ag: 16 – 22 mJ/m 2 , stainless steel: 15 – 60 mJ/m 2 , and Cu: 35 – 78 mJ/m 2 [13, 14, 36]) than those with higher SFEs (such as Al: 135 – 200 mJ/m 2 and Ni: 125 – 300 mJ/m 2 [13, 14]). There are two major types of nanotwins in terms of the formation process: deformation nanotwin (by top-down approach) [37, 38], and growth nanotwin (by bottom-up approach) [39, 40]. Electrodeposition and magnetron sputtering are two popular methods for producing growth nanotwins. Magnetron 17 sputtering was used to synthesize nt Cu samples used in this study. Fig.11 is a schematic showing the deposition process during magnetron sputtering. Figure 11 Schematic of the magnetron sputtering process, adapted from [41]. As shown in Fig.11, the inert gas atoms (Ar in this work) are ionized creating the plasma consisting of positive gas ions (Ar + ) and electrons. The voltage field drives the positive gas ions (Ar + ) to speed towards the target and “knock out” metal atoms. The knocked-out metal atoms are subsequently deposited on the substrate (Si in this work). The magnets in magnetron sputtering help to generate stable plasma with a high density of electrons and positive gas ions near the target, which enhances sputtering efficiency. 18 Figure 12 TEM micrographs of the (a) top surface, (b) cross-section of magnetron sputtered nt Cu [42], and the (c) top surface of electrodeposited nt Cu [43]. Fig.12 (a) and (b) show the top-surface and cross-sectional grain structures of a highly nt Cu foil synthesized by magnetron sputtering. The cross-sectional micrograph [Fig.12 (b)] shows that the Cu has columnar grains aligned parallel to the growth direction, and each columnar grain has a high density of nanotwins perpendicular to the growth direction. The columnar grain size and nanotwin thickness vary with the sputtering conditions: Hodge et al. [42, 44] produced nt Cu with grain size of  700 nm and twin thickness of  40 nm, while the nt Cu of Misra et al. [8, 45, 46] had a grain size of  50 nm and twin thickness of  4 nm. In addition to magnetron sputtering, most electrodeposition studies have produced [29, 43, 47, 48] nt Cu with equiaxed grains, although columnar grained nt Cu [40, 49] was also reported. Fig.12 (c) shows a high 19 density of nanotwins within the equiaxed grains of electrodeposited nt Cu. The typical size of equiaxed grains is between 400 – 600 nm, and the twin thickness can be varied from 4 to 96 nm depending on the electrodeposition parameters [43, 48]. Figure 13 TEM grain orientation mappings of columnar GBs intersected by nanotwins: (a) an alternating columnar GB, and (b) a high-angle columnar GB, indicated by the two black arrows. GB types are as follows: black, high-angle GB; yellow, low-angle GB; red, 3 GB; blue, 9 GB. In (b), the white arrow refers to the growth direction of the foil, and the inset correlates each color with a crystal orientation [50]. The GB character distribution can be altered by the presence of nanotwins. Fig.13 (a) and (b) show the cross-sectional columnar grains of as-sputtered nt Cu; different colors represent different grain orientations as illustrated by the triangle legend at the bottom right. The black arrow of Fig.13 (a) points along a columnar GB shared by 20 two neighboring nt columnar grains. It can be observed that each nanotwin GB (red line) induces a change in the grain orientation, which then leads to an alternating change in the misorientation of columnar GB between high-angle and low-angle. This interesting phenomenon demonstrates that incorporating nanotwins into a metal could largely alter the GB character distribution by adding the low-angle GB segments. Nevertheless, Wang et al. [50] also observed from Fig.13 (b) that a columnar GB (indicated by the black arrow) maintains its high angle even though nanotwins are present. Therefore, while nanotwins could have a big influence on the GB character distribution, how to control it so as to generate a consistent alternating-GB network is still unclear. 2.3.2 Mechanical properties of nanotwinned metals The mechanical properties of nt Cu synthesized by magnetron sputtering have been well-studied, and it is reported [39, 42, 44-46, 51] that nt Cu has an ultra-high strength and stable microstructure under fatigue. Electrodeposited nt Cu samples have ultra-high strength and reasonable ductility due to the presence of the nanotwins [43, 48]. Fig.14 illustrates the stress-strain curve of nt Cu compared with nanocrystalline (nc) Cu and coarse-grained (cg) Cu. Nt Cu has not only much higher tensile strength than nc Cu and cg Cu, but also a decent amount of elongation (13.5%) before fracture. 21 Figure 14 Tensile true stress – true strain curves of as-deposited nt Cu (average grain size, d avg  400 nm) in comparison with nc Cu (d avg  30 nm) and cg Cu (d avg > 100 µm) [43]. Moreover, it is also found that the thickness (or spacing) of nanotwins can have a significant effect on the stress-strain curve of nt Cu. Fig.15 illustrates that while the ductility of nt Cu increases with decreasing twin thickness, the tensile strength reaches its peak value at 15 nm twin thickness. Therefore, a good combination of strength and ductility could be achieved in nt Cu by controlling its twin thickness. Despite the improved overall mechanical properties, the thermal stability and corrosion resistance of nt Cu need to be further studied. 22 Figure 15 Tensile true stress-true strain curves of nt Cu samples (400 nm < d avg < 600 nm) with various twin thicknesses (from 96 nm down to 4 nm) in comparison with twin-free ufg Cu (d avg  500 nm) and cg Cu (d avg  10 µm) [48]. 2.4 Thermal stability Although alteration of the microstructure to improve the mechanical properties is well-studied, the thermal stability of nt metals has received much less attention. Thus, one of the foci in this study is to identify how nanotwins could affect the grain growth process. In this section, we will introduce the fundamentals of thermal stability, and then the relevant publications on the thermal stability of nt metals will be reviewed. 23 2.4.1 The driving force for grain growth Thermal stability evaluates the ability of a metal to maintain its grain structure at elevated temperatures. The grains of a metal always tend to grow at elevated temperatures to reduce GB interface area and lower the total energy, but grain growth can degrade the properties. The driving force of grain growth is energy reduction, and the energy sources for the driving force are: stored deformation energy, GB energy, surface energy, chemical driving force, magnetic field, elastic strain energy, and temperature gradient [52]. Although GB energy is a common driving force during grain growth, different materials would have different major driving forces due to their specific microstructural features. For example, heavily deformed metal would have stored deformation energy as the main driving force due to its high dislocation density. Ufg metals and nc metals may have GB energy as their main driving force due to the presence of large amounts of GBs. The main driving forces for the grain growth of deposited metallic films are GB energy, surface energy, and strain energy. While GB energy reduction would initiate normal grain growth, surface and strain energy reduction could lead to the growth of a specific grain orientation due to the crystal anisotropy [53-58]. Equation 2 shows the relationship between grain growth rate, GB mobility, and driving forces (GB energy, surface energy, and strain energy) [53, 54]: Equation 2 24 Where is the growing grain radius, is the average GB mobility, is the average GB energy, is the average grain radius, is the driving force of surface energy, is the driving force of strain energy, and is the driving force of GB energy. It can be seen from this equation that the grain growth rate is affected by GB mobility and the driving forces. GB mobility depends on GB structure and temperature, and is generally independent on the driving force [52]. At higher temperature, GBs are more mobile. 2.4.2 Grain growth in nanotwinned metals Section 2.3.1 discussed that the grain size of nt metals typically ranges from 50 to 700 nm, which are in the ufg (grain size, 100 nm – 1 µm) and nc (grain size < 100 nm) regime. Although the grains of ufg and nc metals are likely to grow due to their small grain size, incorporating nanotwins into these small grains may help to enhance the thermal stability. Fig.16 shows a comparison of the thermal stabilities between nt Cu and ufg Cu (without nanotwins) at different heat-treatment temperatures. Ufg Cu had a rapid grain growth below 300°C, while the grain growth of nt Cu was modest. Even at 800°C (  0.8 T m ), the grain size of nt Cu remains in the ufg range and shows excellent thermal stability. The cross-sectional TEM micrographs before and after 800°C heat-treatment are shown in Fig.17. It can be observed that although the columnar grain width increased from  40 nm to  540 nm after heat-treatment, the nanotwins in each columnar grain remained. 25 Figure 16 Grain growth at different heat-treatment temperatures for nt Cu (red circles) and ufg Cu (blue triangles). Red and blue lines are the fitted grain growth curves calculated from normal grain growth equation. Adapted from [8]. Figure 17 Cross-sectional TEM micrographs of as-sputtered nt Cu (a) before and (b) after heat- treatment at 800°C for one hour [8, 9]. 26 In addition to the thermal stability studies by Zhang et al., Xu et al. [10] compared the thermal stabilities of electrodeposited Cu with and without nanotwins, and suggested that Cu with nanotwins is more thermally stable than Cu without nanotwins due to strain energy reduction during nanotwin formation. Saldana et al. [11, 12] produced Cu with nanotwins by severe plastic deformation at cryogenic temperature, and they attributed the high thermal stability of nt Cu to its lower vacancy supersaturation and the nanotwin triple junctions. Although these studies have shown that nt metals have higher thermal stability compared with ufg and nc metals, the processing methods and parameters of the various studies are quite different, which results in nt metals with various microstructural features (grain size, grain shape, grain orientation, twin thickness, and GB character distribution). The role of nanotwins in enhancing thermal stability was therefore studied from different perspectives. Consequently, further studies focusing on GB character and grain orientation of nt metals are needed in order to clearly identify how nanotwins may enhance the thermal stability. 2.5 Corrosion behavior Given that corrosion of many industrial metals is prevalent, it is necessary to determine the role of nanotwins in corrosion behavior. In this section, the background knowledge and how GBs influence corrosion behavior are introduced, and the literature on the corrosion resistance of nt metals is reviewed. 27 2.5.1 The role of grain boundary in corrosion resistance Corrosion is the deterioration of materials (mostly for metals) due to their interaction with the environment. The reason of metal corrosion is that corrosion reverts metals back to their original states (minerals and ores) in which they are more stable [59]. However, corroded metals usually have degraded microstructures and properties, which results in the engineering failures of vehicles, bridges, buildings, pipelines, home appliances, water systems, and so on. It is estimated that the cost (> 1.8 trillion USD) of corrosion worldwide equals to 3 – 4 % of the Gross Domestic Product of industrialized countries [60, 61]. To mitigate corrosion, the effect of GBs on corrosion resistance has been studied. There are two major factors of GBs that can affect corrosion resistance: GB character and GB density. Corrosion mainly occurs at high-energy defective area, and studies [62, 63] indicated that GBs with small deviation from the ideal low-energy GB structure would have good corrosion resistance. Specifically, it was reported [62, 64, 65] that low-CSL (  ≤ 29) GBs, due to their ordered structures, have better intrinsic corrosion resistance. They also have improved resistance to solute segregation and precipitation which could lead to sensitization. However, the influence of GB plane orientation on corrosion behavior should also be considered in addition to CSL value. Miyamoto et al. [66] suggested that the corrosion resistance of <110> tilt GBs in Cu is better related to GB plane orientation instead of the CSL value. Fang et al. [65] found that 9 GBs could have different corrosion resistances due to the change in GB plane orientations. 28 Besides GB character, GB density also has a significant impact on corrosion. GB density is generally linked to grain size, as a smaller grain size means a higher GB density, and vice versa. Previous corrosion studies on metals with different grain sizes have shown no uniform conclusion, which suggests that decrease in grain size could have negative, positive, or no effects on corrosion resistance depending on different metals, corrosive environments, and experimental parameters. Some studies have shown that nc metals have lower corrosion resistance as compared with cg metals [67-72]. Fig.18 is a comparison of the polarization curves of nc nickel (d avg  32 nm) and cg nickel (d avg  100 m), and one can see that both the corrosion and passivation current densities of nc nickel are much bigger than those of cg nickel, which indicates that nc nickel has a lower corrosion resistance compared with cg nickel. This decrease in corrosion resistance of nc metal is generally explained by its increased electro-active surface area due to the high density of GBs, and the defects produced during the processing of the nc metal could also increase its corrosion rate [67]. Other studies [73-80], on the contrary, have shown that nc or ufg metals possess better corrosion resistance than cg metals. As illustrated in Fig.19, a nc Cu coating shows lower passivation current density than mc Cu, and the polarization resistance of nc Cu is over two times that of mc Cu [73]. Researchers [73-80] attributed the better corrosion resistance of nc or ufg metals mainly to the homogeneous passive layer formed on these metals due to the fast diffusion in small-grained metals. However, cg metals have much fewer GBs, typically with high concentrations of energy and impurities which results in 29 localized corrosion. Furthermore, there are several other papers indicating that decreased grain size can have little [81, 82] or complex [83-86] effects on the overall corrosion. Therefore, while a structurally ordered GB character is generally believed to be corrosion resistant, there is no unified conclusion about the effect of GB density on corrosion resistance. Figure 18 Potentiodynamic polarization curves of nc (d avg  32 nm) and cg (d avg  100 m) 99.99% nickel using a scan rate of 0.5 mV/s in 1 mol/L H 2 SO 4 at 293K [67]. 30 Figure 19 Potentiodynamic polarization curves of nc (grain size, 56 nm) and mc (grain size, 2 m) Cu coatings using a scan rate of 0.5 mV/s in 0.1 mol/L NaOH solution [73]. 2.5.2 Corrosion behavior of nanotwinned metals Since both GB character and GB density affect the corrosion resistance, it is very interesting to explore how nanotwins, which are bundles of low-energy CTBs, influence corrosion behavior. Meng et al. [15-17] studied the corrosion behavior of nt Ni. Fig.20 compares the curves of nt Ni and cast Ni (presumably without nanotwins) polarized in 0.1 mol/L H 3 BO 3 + 0.025 mol/L Na 2 B 4 O 7 solution, and Table 2 lists the corrosion parameters derived from the polarization curves of nt Ni and cast Ni. It can be generalized from Fig.20 and Table 2 that nt Ni had higher corrosion potential (E corr ), lower corrosion current density (I corr ), lower passivation current density (I pass ), and higher passive layer breakdown potential (E breakdown ) compared with cast Ni, which indicates that 31 nt Ni has a lower corrosion rate and more protective passive layer. The improved corrosion properties were attributed to a thinner and less defective passive layer formed on the top surface of nt Ni [15-17]. In addition, Meng et al. [18] reported that nt Al also had lower corrosion rate and higher pitting resistance than cast Al. Figure 20 Potentiodynamic polarization curves of nt Ni and cast Ni in 0.1 mol/L H 3 BO 3 + 0.025 mol/L Na 2 B 4 O 7 solution [15]. Table 2 Corrosion parameters derived from the polarization curves of nt Ni and cast Ni in Fig.19 [15]. E corr (mV) I corr (mA/cm 2 ) I pass (mA/cm 2 ) E breakdown (mV) nt Ni -345 0.61 0.67 820 cast Ni -391 2.09 4.49 – 23.1 500 32 However, these electrochemical studies [15-18] focused on Ni and Al which have high stacking fault energies and are not likely to form nanotwins, as discussed in Section 2.3.1. Given that these studies [15-18] concentrated on the electrochemical aspects of corrosion rather than microstructural characterization, it is uncertain what fraction of the grains in the Ni and Al samples contained nanotwins. Moreover, how could nanotwins help to form a protective passive layer and improve the corrosion resistance is still an open question. Therefore, it is necessary to fully characterize the microstructures with and without nanotwins, and compare their corrosion behaviors in order to explore the effect of nanotwins on corrosion resistance. This is one of the subjects of the present work. 2.6 Sensitization of Al-Mg alloys 2.6.1 Sensitization and its consequences 5xxx series Al-Mg alloys are widely applied in transportation, pressure vessels, marine, and other applications because of their high strength, good weldability, and favorable corrosion resistance [87]. However, Al-Mg alloys with more than 3 wt% of Mg, such as Al 5083 (4.0 – 4.9 wt% Mg) and Al 5456 (4.7 – 5.5 wt% Mg), are subject to sensitization. This is the precipitation of β phase (Al 3 Mg 2 ) along GBs leading to intergranular corrosion and stress corrosion cracking. Fig.21 shows the Al-Mg phase diagram, and it can be seen that 3 wt% Mg is beyond the solid solution of Mg in α-Al at temperatures below  195 C. The supersaturated Mg diffuses to the GBs or other 33 defective areas in Al-Mg alloys and forms β phase. Studies of Al-Mg sensitization have been primarily focused on Al 5083 (4.0 – 4.9 wt% Mg). These studies observed β precipitation in Al 5083 at elevated-temperature exposure (typically 150 – 200 C) from 1 hour to 40 days [1, 2, 88-90] and lower temperatures (typically 40 – 100C) for longer times ranging from 3 days to 30 months [91-93]. There are fewer sensitization studies on Al 5456 (4.7 – 5.5 wt% Mg), but it was reported [93, 94] that β precipitation in Al 5456 was even more pronounced than in Al 5083 under similar heat-treatment conditions due to the higher Mg content. Figure 21 Al-Mg phase diagram, adapted from [95]. 34 In general, β phase can be continuous or discontinuous along GBs [2, 92]. It has also been found near the pre-existing Mn-rich particles [89-91]. Fig.22 (a) illustrates the continuous β phase along the GB of an Al 5083 sample after heating at 150 C for 189 hours. However, after heating the same sample at 150 C for a longer period of time (333 hours), the continuous β phase observed in Fig.22 (a) broke down and became discontinuous as shown in Fig.22 (b). The thickness of the intergranular β phase typically ranges from 5 to 300 nm [1, 2, 88-92]. Figure 22 Dark field TEM images of (a) continuous and (b) discontinuous intergranular β phase, adapted from [2]. Since β phase is anodic to Al-Mg matrix, it can cause the intergranular corrosion and stress corrosion cracking, and thus deteriorate the mechanical properties of engineering materials [1, 2, 96]. Fig.23 shows the ductility (or normalized strain to failure) of heat-treated Al 5083 in a 3.5% NaCl solution, and the insets present the formation of intergranular β phase. It can be observed that the ductility decreases with 35 increase in heat-treatment time due to the formation of intergranular β phase. At 189 hours, continuous β phase was formed along the GB, and caused the ductility to reach its minimum. For the heat-treatment longer than 189 hours, the ductility recovered as continuous β phase was broken apart into discontinuous ones. Thus, continuous β phase could be more detrimental to ductility, intergranular corrosion, and stress corrosion cracking compared with discontinuous β phase [2]. Figure 23 Normalized strain to failure versus the time of heat-treatment (at 150 C) plot obtained by the constant extension rate testings of Al-5083 samples in a 3.5% NaCl solution. Insets are the TEM images showing β precipitation at selected heat-treatment times. Adapted from [2]. 36 2.6.2 Model of β phase growth The intergranular β phase growth is due to the diffusion of Mg from the grain interior toward GBs. Fig.24 (a) illustrates the Mg concentration profile during β phase growth, and C β , C 0 , and C e are the Mg concentration in β phase, matrix, and equilibrium state respectively. Since β phase (Al 3 Mg 2 ) is rich in Mg (37 wt%), the Mg content in its adjacent matrix area is depleted and transferred to the β phase in order to support is growth. The growth rate (v) under the model of Fig.24 (a) is [97]: Equation 3 Where D is the interdiffusion coefficient, and dC/dx is the concentration gradient at the interface between Al-Mg matrix and the growing β phase. The model in Fig.24 (b) simplifies the Mg diffusion profile of Fig.24 (a) into a line with L as the diffusion zone width and (C 0 - C e ) as the Mg concentration difference. The two shaded areas in Fig.24 (b) have to be equal due to the conservation of Mg content [97]: (C β – C 0 ) ∙ x = L ∙ (C 0 - C e )/2. Since the “x” in this equation indicates the thickness of β phase on one side of the GB, the total thickness (T) is two times of “x” and can be derived as: Equation 4 Equation 4 shows the relationship between β phase thickness and diffusion zone width. Combining Equation 3 and Equation 4, and knowing that dC/dx = (C 0 -C e )/L under the simplified model, the relationship between β phase growth rate (v) and β phase thickness (T) can be obtained as below: 37 Equation 5 Figure 24 The (a) model and (b) simplified model (originated by Zener [98]) of solute concentration profile during diffusion controlled β phase growth [97]. C β , C 0 , and C e refer to the Mg concentrations in β phase, Al-Mg matrix, and equilibrium state. While Equation 4 correlates β phase thickness (T) with diffusion zone width (L), Equation 5 presents the inversely proportional relationship between β phase growth rate (v) and β phase thickness (T). These two equations will be used in this study to further analyze the effects of grain size, dislocations, and GB planes on β phase growth during sensitization. 38 2.6.3 Grain boundary effects While many studies demonstrate the dependence of β precipitation on the Mg content [94], alloying elements (e.g. Zn) [88], welding [99], and elevated-temperature exposure [100], there are very limited studies about the effect of GB character on β precipitation. A GB has 5 degrees of freedom: 3 degrees define GB misorientation, and the other 2 degrees define GB plane orientation. From the perspective of GB misorientation, Davenport et al. [22] used 10% H 3 PO 4 to observe the β phase of a sensitized Al 5182 sample, and found that low-angle GBs (≤ 15 ) were not etched and thus were highly resistant to β precipitation as shown in Fig.25. In contrast, Kaigorodova et al. [23] observed β phase along low-angle GBs (5 - 10), and Scotto D’Antuono et al. [24] further suggested that β precipitation along low-angle GBs is more prevalent than that on high-angle GBs. Figure 25 The relationship between GB misorientation angle and attacked mode obtained from etching a Al 5182 sample sensitized at 150 C for 10 hours [22]. 39 Nevertheless, these previous studies [22-24] agreed that in addition to GB misorientation, GB plane orientation could largely affect the sensitization behavior and therefore should be further explored. Evidence of this can be observed in Fig.25: in the range of misorientation angle (θ) bigger than 25 , GBs could exhibit different sensitization behaviors even if they have the same GB misorienation. This indicates that GB misorientation is not the only factor influencing β precipitation, but GB plane orientation also needs to be considered. To date, the role of different GB plane orientations in β precipitation has not been studied. However, there is one study [101] that examined the interface between the β phase and the Al-Mg matrix. Yuan [101] showed in Fig.26 (a) that the morphology of intergranular β phase varies, and as indicated by the box “A”, β phase preferentially grew along certain crystallographic directions. Yuan [20] analyzed the normal directions (N 1 , N 2 ) of the two β phase interfacial plane as shown in Fig.26 (b). After surveying six β phase sites, Yuan [20] concluded that β phase grew primarily on the {111} interfacial plane with {211} as a secondary plane. It was further observed that interface planes with larger divergence angles from {111} orientation had improved resistance to β precipitation [20]. Nevertheless, the study focused on the dependence of β phase growth on interface plane orientation instead of GB plane orientation, and more studies need to be carried out to directly relate GB plane orientation with β precipitation behavior. 40 Figure 26 (a) Bright-field image of an intergranular β phase in a sensitized Al 5182; (b) magnified image of a β phase with the normal directions (N 1 , N 2 ) of its interfacial planes. Adapted from [101]. 41 Chapter 3 Characterization and test methods Test methods, such as heat-treatments, polarization, immersion corrosion, and chemical etching, were used in this work in order to evaluate the thermal, corrosion, and sensitization properties of the samples. Before and after each test, the samples were characterized by scanning electron microscope (SEM), X-ray diffraction (XRD), EBSD, focused ion beam (FIB), and TEM, so that the surface topography, grain orientation, GB character, grain size, and chemical composition can be revealed. The details of these characterization and test methods are described in this chapter. 3.1 Characterization methods 3.1.1 X-ray diffraction (XRD) XRD was used in this study to detect the surface grain orientation and chemical compounds at a macroscopic scale (  10 mm). X-rays are electromagnetic waves with wavelength (λ) in the order of 0.1 nm, which is much shorter than visible light [102]. During XRD scan, the incident X-ray beams are deflected by the crystal planes on the surface of the material as shown in Fig.27. Prior to interaction with the material, beams 1 and 2 are in-phase; but after deflected by the crystal planes, they will not be in-phase unless their path difference (SQ + QT) is equal to one or multiple times of the wavelength (nλ) as shown in Fig.27. Since SQ = QT = d hkl ∙ sinθ, the Bragg’s law can be derived as following: 42 Equation 6 Where n is an integer, λ is the wavelength of X-ray, d hkl is the spacing of the {hkl} atomic plane, and θ is the angle of incident beams as shown in Fig. 27. From XRD scan results, we can obtain the θ value of each peak where there is a constructive (or in- phase) interference of the deflected X-ray beams. Therefore, the atomic plane spacing (d) can be calculated by Equation 6. With d known, the Miller indices {h k l} of a crystallographic plane (or grain orientation) can be obtained from the following equation: Equation 7 Where “a” is the lattice parameter. Figure 27 Schematic of Bragg’s diffraction by crystal planes [103]. 43 In addition to grain orientation detection, XRD was also used to identify the passive layer compounds of Cu samples after corrosion. Table 3 lists the d hkl and 2θ of typical grain orientations {h k l} of Cu, CuCl, Cu 2 O, and Al. Table 3 Typical grain orientations{hkl} and the corresponding d and 2θ of Cu (PDF#98-000- 0172), CuCl (PDF#98-000-0323), Cu 2 O (PDF#98-000-0186), and Al (PDF#00-004-0787). {hkl} d (Å) 2θ ( ) {hkl} d (Å) 2θ ( ) Cu {111} 2.0871 43.317 CuCl {111} 3.1281 28.512 {200} 1.8075 50.449 {220} 1.9156 47.422 {220} 1.2781 74.126 {311} 1.6336 56.268 Cu 2 O {111} 2.4647 36.424 Al {111} 2.3380 38.473 {200} 2.1345 42.308 {200} 2.0240 44.739 {220} 1.5093 61.376 {311} 1.2210 78.229 3.1.2 Scanning electron microscope (SEM) Fig.28 shows the JSM-7001F-LV field-emission SEM that was frequently used in this study to examine the surface topography, chemical composition, and grain orientation of materials. Once a SEM electron beam interacts with the specimen surface, secondary electrons, backscattered electrons, and characteristic X-rays are generated at different depths as shown in Fig.29. Secondary electrons are the products of inelastic scattering where the incident electrons transfer the kinetic energy to the specimen atoms and eject the electrons of the specimen atoms out of their orbitals to become secondary electrons [102, 104]. The energy of a secondary electron is as low as several to tens of keV, so they can only escape within a shallow depth (5 – 50 nm) from the specimen 44 surface as illustrated in Fig.29 [102, 104]. Different from the secondary electrons, backscattered electrons are the incident electrons elastically scattered by the specimen atoms, so they can retain 60 – 80% of the incident electron energy [102, 104]. Fig.29 shows that they can escape from a deeper level (50 – 300 nm) under the specimen surface. Backscattered electrons can reflect the elemental and crystallographic information [102, 104], and will be further discussed in Section 3.1.3. On the other hand, secondary electrons are the primary signal for topographic contrast, so secondary electron imaging was used to observe the surface topographies of corroded Cu and etched Al-Mg alloys. Figure 28 JSM-7001F-LV field-emission SEM equipped with an EBSD analytical detector. In addition to SEM topography, this study also used the energy dispersive spectroscopy (EDS) under SEM to detect the elemental compositions of materials. 45 Fig.29 illustrates that characteristic X-rays can also be generated by the interaction between the electron beam and surface atoms. Since the characteristic X-ray energy of an element is the energy difference between its outer-shell and inner-shell electrons, EDS can determine which element it is from the characteristic energy [102, 104]. The energy of characteristic X-ray is typically bigger than those of secondary and backscattered electrons, thus characteristic X-ray can escape from a deeper depth as shown in Fig.29. The resolution of EDS is typically in the range of 0.5 – 5 µm [102, 105, 106]. Figure 29 The electron – specimen atom interaction volume under the specimen surface [102]. 46 3.1.3 Electron backscatter diffraction (EBSD) Figure 30 Schematic set-up of a conventional EBSD system [107]. EBSD tests were carried out by the SEM in Fig.28 with an EDAX Hikari detector to characterize the GBs, grain orientations, and grain sizes of the materials. Conventional EBSD mainly consists of three parts as illustrated in Fig.30: SEM which generates the electron beam, the 70°-tilted polished sample, and the image acquisition and processing device. In Fig.30, the electron beam goes down and interacts with the 70° tilted sample, and then a backscattered Kikuchi diffraction pattern can be formed on the phosphor screen. Fig.31 (a) shows the interaction between the electron beam and the crystallographic planes of the specimen. After the electron beam strikes the specimen surface, electrons will be scattered by the specimen atoms in all directions. Some scattered electrons, which travel in the directions that satisfy the Bragg’s law (Equation 6) of certain crystallographic planes, will then be channeled to form the Kikuchi lines and 47 pattern as shown in Fig.31 (b). Since Kikuchi patterns bear the information of crystallographic plane orientations, grain (or crystal) orientations can be obtained by software indexing of these patterns. Fig.32 shows a schematic grain orientation map: different colors represent different crystal orientations, and the area with similar color forms a grain. Figure 31 (a) Schematic of the interaction between electron beam and tilted specimen, and the subsequent Kikuchi line generation; (b) A Kikuchi pattern from cadmium at 20 keV [108]. Figure 32 Schematic of EBSD grain orientation image. Grain orientations are presented by the colored crystal lattices [109]. 48 Besides grain orientation map, another important application of EBSD is GB characterization. Fig.33 shows a typical EBSD GB character distribution map of a GBE brass, and it can be seen that the GBE brass contains a large fraction of 3 GBs (red lines) [5]. In addition to the GB misorientation which contains three out of the five GB parameters, one additional GB parameter (GB trace on the surface) can also be detected by a single-section EBSD study. With this additional GB parameter, EBSD can differentiate most incoherent 3 GBs from coherent 3 GBs by GB trace analysis [110, 111]. Figure 33 GB character distribution of a GBE brass. Black, red, and blue lines represent ordinary, 3, and 9 GBs respectively. Adapted from [5]. Although it is fairly easy for conventional EBSD to analyze the grain size and GB misorientation of a well-polished crystalline sample, the optimum spatial resolution of it is only  20 nm and most times greater than 20 nm [109]. So it is difficult for 49 conventional EBSD to detect nanocrystalline grains (grain size < 100 nm). Nevertheless, a new technique called “transmission EBSD” was recently initiated by Keller and Geiss [112]. Different from conventional EBSD, transmission EBSD collects the transmitted diffraction patterns from a thin film, and can achieve a much higher resolution (< 10 nm) [112-114]. Figure 34 The experimental set-up of transmission EBSD. The dark blue, red, and light blue lines refer to the incident electron beam, plane of the thin film, and the transmitted electrons towards the EBSD detector [112]. The experimental set-up of transmission EBSD is shown in Fig.34. TEM thin film (typically thickness < 100 nm) is used as the specimen, so that the incident electrons can transmit through the thin film and be collected by the EBSD detector (see the light blue lines in Fig.34). As discussed earlier, the mechanism of conventional EBSD is that 50 the incident electrons are scattered in all direction and some of them are channeled to generate the diffraction pattern which bears the crystal orientation information; in the case of transmission EBSD, the difference is that the transmitted electrons instead of reflected electrons generate the diffraction patterns, but the general mechanism is similar to conventional EBSD. Additionally, compared with the conventional EBSD as shown in Fig.30, transmission EBSD places the thin film higher relative to the EBSD detector and tilts it in the opposite direction of the EBSD detector, so that the transmitted diffraction pattern can be captured. The uniqueness of transmission EBSD is that the “thin” film largely reduces the interaction volume (shown in Fig.29) generated by the incident electrons and specimen atoms, so the spatial resolution of transmission EBSD can be much smaller than that of conventional EBSD. In this study, while the conventional EBSD was widely applied to characterize the microstructures of cg (grain size ≥ 1 µm) and ufg (100 nm < grain size < 1 µm) samples, transmission EBSD was used to characterize nc (grain size < 100 nm) samples. 3.1.4 Focused ion beam (FIB) JEOL JIB-4500 FIB was applied in this study for the purpose of cross-sectional grain imaging and TEM sample preparation. Different from SEM, the incident beam of FIB is an ion beam. 51 Figure 35 Schematic illustration of the interaction between an incident Ga + ion and crystal lattice atoms. R p and R l are the projected and lateral ranges of the implanted Ga + ion [115]. Fig.35 illustrates the interaction between an incident Ga + ion and the specimen atoms. As the Ga + ion penetrates the specimen surface, its kinetic energy can be 52 transferred to the electrons of the specimen atoms during collision leading to the emitted secondary electrons. Moreover, if the transferred energy exceeds a critical value, the atom itself can be knocked out of its original position creating an interstitial-vacancy pair as shown in Fig.35 [115]. When the collision happens near the specimen surface, the atoms could be knocked out of the specimen and become sputtered atoms. Although both SEM and FIB can use secondary electrons to image, clear grain contrast can only be obtained by FIB. Fig.36 (a) and (b) present the interaction of a Ga + ion with two distinctly oriented grains. In (a), the grain is favorably oriented so that the Ga + ion can channel through the atoms without much interaction, thus fewer secondary electrons are emitted from the surface; in contrast, the grain in (b) is unfavorably oriented for ion channeling, and there are active interactions between the Ga + ion and the lattice atoms which results in more emitted secondary electrons as shown by the red arrows [104, 115]. This mechanism leads to the strong grain contrast in the secondary electron images of FIB which was used in this study to detect grain size and morphology. In addition, the amount of emitted secondary electrons depends on the atomic mass. Fig.36 (c) shows that heavier atoms can generate more secondary electrons during the interaction with the ion compared with lighter atoms (a) and (b). In this study, it was observed that the FIB images of Cu samples showed a clear grain contrast, but those of Al-Mg samples had a blurred grain contrast mainly due to their smaller atomic mass. One thing to note is that the samples were tilted 52  during cross-sectional FIB imaging, so the actual length in the vertical direction is  1.27 times the vertical length shown in the cross-sectional FIB micrographs. 53 Besides grain imaging, FIB was also used to prepare TEM samples. Fig.35 demonstrates that some atoms of the specimen can be sputtered away due to the collision with the incident ions. Using ion beam sputtering, we milled away the surrounding materials, and used a probe tip to extract the area of interest out of the sample as shown in Fig.37 (a). It should be noted that in both FIB imaging and lift-out, a carbon layer was pre-deposited on top of the interested area to protect it from Ga + ion damage. After the lift out, the area of interest was then attached to the TEM grid and subsequently detached from the probe tip in Fig.37 (b). The area of interest was then further thinned until its thickness is below  100 nm which is suitable for TEM observation. Figure 36 Schematic images illustrating the effect of (a), (b) grain orientation, and (c) atomic mass on the emitted secondary electrons. The orange atoms in (c) have bigger atomic mass compared with those in (a), (b). Adapted from [115]. 54 Figure 37 (a) The lift-out of an area of interest from the sample by a probe tip; (b) The attachment of the area of interest to a TEM grid for thinning and TEM observation. 3.1.5 Transmission electron microscope (TEM) TEM was mainly used in this study to observe the nc grains and β phase of Al-Mg samples. The formation process of a TEM image by diffraction contrast is shown in Fig.38 (a). The electron beam interacts with the sample according to the Bragg’s condition (Fig.27), and most electrons will either transmit or be diffracted depending on different grain orientations. In Fig.38 (a), the diffracted beam (I d ) forms a diffraction pattern on the back focal plane, while the transmitted beam (I t ) travels further and forms a bright-field image on the image plane. Fig.38 (b) illustrates the details of the bright-field image formation: the objective aperture selects the transmitted beam (I t ) to go through, but the diffracted beam (I d ) is blocked; on the other hand, for dark-field image formation, the objective aperture only allows the pass of one diffracted beam as shown in Fig.38 (c). However, some grains may not be observable in the TEM images formed by diffraction contrast, because only a fraction of the grains could satisfy the Bragg’s condition and be revealed [102]. 55 Figure 38 TEM image formation by diffraction contrast: (a) detailed formation process; (b) bright-field image formation; (c) dark-field image formation. Adapted from [102]. Besides TEM imaging, the EDS and electron energy loss spectroscopy (EELS) were applied to detect the chemical compositions of sensitized Al-Mg samples under the scanning-TEM (STEM) mode. The mechanism of characteristic X-ray formation in STEM-EDS is the same as SEM-EDS discussed in Section 3.1.2. However, since STEM-EDS uses thin foil (typical thickness < 100 nm), it largely reduces the interaction volume and thus improves the resolution. The spatial resolution of STEM-EDS can be less than 2 nm [116]. Different from EDS which collects characteristic X-rays, EELS collects the energy loss of the transmitted electrons. Fig.39 (a) illustrates the process of inelastic scattering between an incident electron and an atom [117]. After scattering, the energy of 56 the incident electron decreases by ∆E and becomes E 0 - ∆E. This ∆E was transformed to the inner-shell electron (energy, E kin ) of the atom leading to its ejection out of the atom. The outer-shell electron of the atom then fills the inner-shell vacancy, and may emit a characteristic X-ray (energy, hυ). This process demonstrates that both the characteristic X-ray (energy, hυ) and the loss energy (∆E) of the incident electron contain the information relating to the energy difference between orbital shells of the atom [104, 117]. Therefore, the elemental composition can be characterized by both characteristic X-rays (EDS) and the energy loss (EELS). Figure 39 (a) Inelastic scattering of the incident electron on the specimen atom; (b) EELS energy loss spectrum of boron nitride (BN). Adapted from [117]. Fig.39 (b) shows the high energy loss region (∆E > 50 eV) of an EELS energy loss spectrum. The two major peaks in the spectrum represent the critical energies needed to eject the inner-shell electrons, so they can be used to evaluate the chemical composition of the material [104, 117]. In Fig.39 (b), the peaks refer to B and N atoms, 57 respectively. The spatial resolution of EELS is below 1 nm, generally higher than EDS [116]. 3.2 Test methods 3.2.1 Heat-treatments Heat-treatments were carried out in thermal, corrosion and sensitization studies. All samples were heat-treated under vacuum environment in order to prevent oxidation. Fig.40 shows the vacuum furnace system used in this study: the lower part is the vacuum pump, and the upper part is the furnace. The vacuum level largely depends on the heat- treatment temperature and holding time, and is generally better than 3 × 10 -5 Torr. After heat-treatments, no oxide was observed on the samples. Figure 40 MTI-GSL1100X vacuum furnace system. 58 Fig.41 presents the temperature profiles during heat-treatments. The heating rate was 10 C/min, and all samples were cooled in the furnace after heat-treatments. For the thermal studies of nt Cu samples, the holding time was three hours as shown in Fig.41, whereas the holding time in the sensitization study of Al-Mg samples was one week. Figure 41 Temperature profiles of the heat-treating, holding and cooling processes used in this study. 3.2.2 Polarization Electrochemical polarization was used to evaluate the corrosion properties of nt Cu samples. Fig.42 (a) and (b) shows the actual and schematic electrochemical cell. The cell is mainly consisted of the electrolytic solution and three immersed electrodes: 59 working electrode, reference electrode, and counter electrode. The beaker in Fig.42 (a) contained 150 ml artificial seawater which is a naturally aerated 3.5% (35 g/L) NaCl solution prepared by mixing ultra-pure water with NaCl. NaOH was also added to the solution to adjust its pH value to 8.0  0.1. The temperature of the NaCl solution was maintained at 21.5  1 C. Figure 42 The (a) actual and (b) schematic electrochemical cell showing the three electrodes (black clip – working electrode; yellow clip – counter electrode). (c) Epoxy-mounted sample as the working electrode, and (d) the Gamry Reference 3000 potentiostat. For the three electrodes, the working electrode of the electrochemical cell was Cu foils with different nanotwin fractions. The Cu foil was covered by epoxy as shown in Fig.42 (c), so that only a specific area of the surface can be exposed to the electrolytic 60 solution. The reference electrode is a silver-silver chloride (Ag/AgCl) reference electrode, and its conversion with saturated calomel electrode (SCE) and standard hydrogen electrode (SHE) is presented in Table 4. Platinum wire was used as the counter electrode in this study. Fig.42 (d) presents the potentiostat which controlled the potential (between working and reference electrodes) and measured the current (between working and counter electrodes) during electrochemical tests. Table 4 Conversion of potentials among Ag/AgCl, SCE, and SHE. Reference Electrode E vs. SCE E vs. SHE Silver/Silver Chloride (Ag/AgCl) E Ag/AgCl ‒ 0.042 E Ag/AgCl + 0.198 There are two types of polarization methods used in this study: linear polarization and potentiodynamic polarization. Linear polarization was applied to evaluate the polarization resistance (R p ) of Cu foils with different nanotwin fractions. Given that the potential and current density have a linear relationship near the corrosion potential (E corr ), linear polarization does a short-range potential sweep typically from - 20 to + 20 mV relative to the corrosion potential (E corr ). Fig.43 illustrates the linear relationship between potential and current density, and the slope of this curve is the polarization resistance (R p ). R p can be related to corrosion current density by the Stern-Geary equation [118, 119]: Equation 8 61 Where I corr is corrosion current density (directly related to corrosion rate), R p is polarization resistance, and β a and β c are anodic and cathodic Tafel slopes that will be further discussed in the potentiodynamic polarization. Figure 43 Schematic plot of a linear polarization curve [120]. In addition to linear polarization, potentiodynamic polarization was used in order to evaluate the corrosion current density and passivation property of Cu foils. Fig.44 shows the schematic of a potentiodynamic polarization curve generated by a potential 62 scan typically from negative hundreds of millivolts to positive hundreds of millivolts relative to the corrosion potential. Different from linear polarization, potentiodynamic polarization curve has the logarithm of current density as the horizontal axis. Figure 44 Schematic of the potentiodynamic polarization curve for a passive material. “C. D.” stands for current density. Adapted from [120]. The curve in Fig.44 can be divided into two regions: the region with potential below the primary passivation potential (E pp ) is the Tafel region, and the region with potential above E pp is the passive region where a passive layer is formed on the material. The Tafel region is typically from – 250 to + 250 mV relative to the corrosion potential 63 (E corr ). From the Tafel region, the slopes (β a and β c ) of the anodic and cathodic linear Tafel branches can be measured as marked in Fig.44. Therefore, given that R p is known from the linear polarization curve, the corrosion current density (I corr ) could be calculated through Equation 8. While this method was used in this study to calculate the corrosion current density of the Cu samples, there is an alternative way that can directly measure the corrosion current density. In the Tafel region, the anodic and cathodic linear branches could be extrapolated as shown by the dashed lines of Fig.44, and they intersect at a point. The (x, y) values of this intersection point can be measured, and they are corrosion current density and corrosion potential, respectively [120]. Above the Tafel region, Fig.44 shows that the current density reaches its critical value (I cc ) at E pp , and then decreases dramatically due to the formation of a protective passive layer on the sample. The region above E pp is thus the passive region. When the potential increases to the passivation point marked in Fig.44, the current density reaches a minimum value named passivation current density (I pass ). I pass reflects the protectiveness of the passive layer, and a lower I pass indicates a more protective passive layer. With further increasing in the potential, there is a break-down of the passive layer resulting in the increasing current density as shown in Fig.44. For some materials, there could be a secondary passive region and subsequent break-down when the potential is even higher. After each polarization scan, the sample was immediately taken out of the solution to prevent further corrosion, and gently rinsed in order to remove the residue NaCl solution on the surface. 64 In sum, the corrosion rate and passivation behavior of a metal can be evaluated by linear polarization and potentiodynamic polarization. Thus, both polarization tests were carried out in this thesis. 3.2.3 Immersion corrosion In addition to the short-term polarization, long-term immersion corrosion was performed on Cu samples in order to further analyze their passivation behaviors. The electrolytic solution of immersion corrosion tests was 3.5% NaCl with pH  8.0, which is the same as that used in polarization tests. Each sample was placed in a close-lid container with 100 ml electrolytic solution for a total of 10 days. Since the pH level and ion concentration of the solution may change due to corrosion reactions, the NaCl solution in each container was replenished every 2 days. After immersion corrosion, the samples were gently rinsed to remove the residue NaCl solution on the surface. 3.2.4 Chemical etching The purpose of chemical etching was to etch away the β phase and analyze the extent of β precipitation in the Al-Mg alloy. Phosphoric acid (H 3 PO 4 ) is generally accepted [22, 93, 100, 121-124] as an effective etchant to reveal the β phase; the etching conditions from previous studies, including H 3 PO 4 concentration, etching temperature and time, are listed in Table 5. Zhu [93] further suggested that H 3 PO 4 may even etch away material which did not contain β phase and thus over-etch the sample. Therefore, it 65 is crucial to select an optimum etching condition which not only etches the β phase but also does not affect the areas without β phase. Table 5 β phase etching conditions. H 3 PO 4 concentration a Etching temperature ( C) Etching time (min) Zhu [93] 10% 50 2 Popovic et al. [121] 10% 50 2.5 Jain et al. [122] 10% 55 0.5 Davenport et al. [22] 10% 60 0.5 Gunson [123] 10% 60 1.5 Oguocha et al. [100] 10% 60 5 Unocic et al. [124] 40% 35 3 a The concentration of H 3 PO 4 is in volume percent. In order to search for an etching condition that does not over-etch the sample, the as-received Al 5456 with few β phase was selected. The etching condition “10% H 3 PO 4 at 50 C for 2min” was initially attempted, and Fig.45 shows the etched surface of the as- received Al 5456. Both the optical (a) and SEM (b) micrographs show the apparent intergranular etching of the as-received Al 5456, and the EBSD (c) map further confirms that the GBs were preferentially etched. However, the as-received Al 5456, which had few intergranular β phase, was not supposed to show intergranular etching under an optimum etching condition. So the initial etching condition “10% H 3 PO 4 at 50 C for 2min” over-etched the sample and was modified. 66 Figure 45 The (a) optical, (b) SEM, and (c) EBSD surface micrographs of the as-received Al 5456 after 10% H 3 PO 4 etching at 50 C for 2 minutes. The triangle legend correlates the colors in (c) with specific grain orientations. Subsequently, a series of etching conditions with lower etching temperature and shorter etching time were tried, and “10% H 3 PO 4 at 35 C for 1 min” was found to be effective in revealing β phase without over-etching. Fig.46 shows the optical graphs of as-received Al 5456 (with few β phase) and sensitized Al 5456 (heat-treated at 175 C for 7 days to induce β precipitation). Before etching, Fig.46 (a) and (b) present well-polished surfaces with the pre-existing Al-Fe-Mn and Mg-Si rich precipitates (ref. [100]) of as- received Al 5456 and sensitized Al 5456, respectively. After etching in 10% H 3 PO 4 at 35C for 1 min, Fig.46 (c) shows that the surface of as-received Al 5456 does not change much as compared with that before etching, which indicates that this etching condition 67 did not over-etch the GBs (without the presence of β phase). In comparison, for sensitized Al 5456 which contained β phase after the prolonged heat-treatment, its surface [Fig.46 (d)] was clearly etched compared with the other figures. This suggests that the current etching condition could also reveal the existing β phase. Therefore, “10% H 3 PO 4 at 35 C for 1min” was selected as the etching condition in this study to evaluate β precipitation. Figure 46 The optical surface graphs of as-received Al 5456 and sensitized Al 5456 before (a, b) and after (c, d) etching in 10% H 3 PO 4 at 35 C for 1 min. 68 Chapter 4 The effect of nanotwins on the thermal stability of Cu This study evaluated the thermal stability of nt Cu as compared with ufg Cu in the perspective of GB character and grain orientation, and further explored how nanotwins could affect the thermal behavior. Ufg Cu and nt Cu samples were heated at regular intervals until the grain structure of nt Cu was completely changed. EBSD, FIB, and XRD were used to characterize the GB migration and texture change before and after each heat-treatment. The major part of this thermal study was published in the Journal of Materials Research under the title “Thermal stability of highly nanotwinned copper: The role of grain boundaries and texture” by Y. Zhao, T.A. Furnish, M.E. Kassner, and A.M. Hodge [125]. 4.1 Experimental procedures Ufg Cu and nt Cu foils were synthesized by magnetron sputtering using high- purity Cu (99.999%) with different sputtering parameters. Both foils were “freely” removed from the substrate and were handled as flat and undistorted free-standing foils. The thickness of ufg Cu foil was  25 m, and that of nt Cu foil was  170 m; nt Cu was polished down to  65 m in order to prepare for EBSD analysis. As shown in Fig.47, fiducial marks were made on the samples by nanoindentation to make sure that the exact same area can be scanned by EBSD before and after each heat-treatment. The ufg Cu and nt Cu freestanding foils were heat-treated sequentially at 200, 300, and 400°C for 3 hours under vacuum (1 × 10 -4 – 1 × 10 -5 Torr); nt Cu was further heat-treated to 69 600°C for 3 hours. XRD was carried out on the top surfaces of the samples to analyze the macrostructural texture and ensure that no copper oxides were formed after heat- treatments. FIB imaging was done on both top and cross-sectional surfaces of ufg and nt Cu foils to characterize the grain morphology and GBs in the as-prepared condition and after 400°C heat-treatment; top-surface FIB grain imaging was also done on nt Cu after 600°C heat-treatment. EBSD was used to analyze top-surface grain size, GB character, and grain orientation. Since the nanotwin thickness is  40 nm, EBSD was not performed on the cross-section due to resolution limitations. Figure 47 EBSD scan area and its near-by fiducial nanoindents in the thermal study. The initial EBSD scan area of the as-prepared samples was 10 m × 10 m. Because of the grain growth at elevated temperature, the scan area of ufg Cu increased to 60 m × 60 m after heat-treatment above 200°C, and that of nt Cu increased to 180 m × 180 m after heat-treatment at 400°C. After the EBSD scans, TSL OIM Analysis 6.1 70 software was used to analyze the result based on these scan areas. For grain size measurement, grains were defined by GBs with misorientation bigger than 5°, and 3 and 9 GBs were excluded from grain size measurement [126, 127]. Moreover, 3 GBs, low-angle GBs (5 - 15°), and high-angle GBs (15 - 63°) were analyzed. In order to differentiate coherent GBs from incoherent GBs, GBs were reconstructed by the software, and boundary trace analysis [110, 111] was applied to examine if the GB plane could coincide with the twinning plane. The length fractions of 3 and coherent 3 GBs were also calculated. 4.2 As-prepared samples Fig.48 (a) and (b) show the top surface and cross-sectional FIB micrographs of ufg Cu, respectively. The grain size is calculated to be  320 nm on the top-surface, and the cross-section presents columnar grains with an average grain size of  400 nm. Fig.48 (c) and (d) show the top-surface and cross-sectional FIB micrographs of nt Cu: the top surface grain size is  700 nm, and a high density of growth nanotwins (twin thickness,  40 nm) can be observed in the cross-section. These nanotwins align normal to the growth direction indicated by the arrow in Fig.48 (d), and are parallel to the top surface of the foil. The above microstructural values are consistent with the results of earlier works using TEM [42, 128]. From the observation of FIB micrographs, the microstructural differences between ufg Cu and nt Cu are primarily the nanotwins and columnar grain size. 71 Figure 48 FIB micrographs of as-prepared ufg Cu on (a) top surface, (b) cross-section, and nt Cu on (c) top surface, (d) cross-section. The arrows in (b) and (d) indicate the growth direction of foils. 4.3 Grain growth at elevated temperature Before and after each heat-treatment, the grain growths of ufg Cu and nt Cu were evaluated. Fig.49 shows the top-surface grain morphologies of ufg Cu in the as-prepared condition and after each heat-treatment; note that each color represent a grain, not a grain orientation. Fig.49 (a) presents the grains of the as-prepared ufg Cu, which is consistent with the FIB result [Fig.48 (a)]. After 200°C heat-treatment, a rapid grain growth can be observed in Fig.49 (b), and the average grain size increased from  330 nm (as-prepared) to  960 nm. A large amount of ultra-fine grains were consumed by several newly formed big grains (5 – 30 m in diameter). At 300 and 400°C, these big grains 72 continually consumed the remaining ultra-fine grains, but the size and shape of them stayed stable as shown in Fig.49 (c) and (d). Figure 49 EBSD top-surface grain morphologies of ufg Cu (a) as-prepared, and after (b) 200°C, (c) 300°C, (d) 400°C heat-treatments. Each color represents a grain, not a grain orientation. Fig.50 (a) shows the top-surface grain morphology of as-prepared nt Cu, which is consistent with the FIB result [Fig.48 (c)]. The grains stayed stable at 200°C in Fig.50 (b), and did not appear to grow even at 300°C in Fig.50 (c). At 400°C [Fig.50 (d)], however, several very big grains can be observed. 73 Figure 50 EBSD top-surface grain morphologies of nt Cu (a) as-prepared, and after (b) 200°C, (c) 300°C, (d) 400°C heat-treatments. Each color represents a grain, not a grain orientation. The quantitative data of grain size evolution after each heat-treatment is illustrated in Fig.51. For nt Cu, Fig.51 (a1), (a2), and (a3) confirms that most grains stayed stable with only a little increase in grain size from  700 nm in the as-prepared condition to  750 nm after 300°C heat-treatment. After heat-treatment at 400°C, while most grains of nt Cu had a slight increase in grains size, very big grains with size ranging from 6 to 70 m formed as shown in the inset of Fig.51 (a4), and the average grain size of nt Cu reached  1200 nm. This is in contrast to ufg Cu which had a rapid grain growth after heat-treatment at 200°C as observed in Fig.51 (b1) and (b2). In Fig.51 (b3) and 74 (b4), there were sharp drops in the number of small grains (≤ 5 m in diameter), which indicate that the remaining small grains were gradually consumed by the newly formed big grains, and this is consistent with the continual increase in average grain size. Figure 51 EBSD top-surface grain size distribution of nt Cu (a1) as-prepared (average grain size d = 697 nm), and after (a2) 200°C (d = 734 nm), (a3) 300°C (d = 754 nm), (a4) 400°C (d = 1195 nm) heat-treatments; and those of ufg Cu (b1) as-prepared (d = 328 nm), and after (b2) 200°C (d = 963 nm), (b3) 300°C (d = 1871 nm), (b4) 400°C (d = 4364 nm) heat-treatments. The insets represent size distributions of bigger grains. 75 To further analyze the extent of grain growth, the top-surface and cross-sectional FIB micrographs of ufg Cu and nt Cu after 400°C heat-treatment were obtained and shown in Fig.52. For ufg Cu [Fig.52 (a) and (b)], many annealing twins can be observed on both the top-surface and cross-section; for nt Cu [Fig.52 (c) and (d)], one can see that several very large grains formed among many small grains on the top surface, and these large grains are without nanotwins as observed from the cross-sectional view. At this point, it is still not clear how these nanotwins and annealing twins affect the thermal stability of ufg Cu and nt Cu, and this will be further discussed in the next section. Figure 52 FIB micrographs after 400°C heat-treatment of ufg Cu on (a) top surface, (b) cross- section, and nt Cu on (c) top surface, (d) cross-section. The arrows in (b) and (d) indicate the growth direction of foils. 76 4.4 Grain boundary migration Before discussing GB migration, one must take into account the sample purity. Since both ufg Cu and nt Cu foils were of high purity (99.999%), there are few impurities to pin GBs, so that GBs can easily migrate [129]. Fig.53 illustrates the EBSD top-surface GB distributions of ufg Cu and nt Cu in the as-prepared condition and after 400°C heat- treatment; 3 GBs, low-angle GBs, and high-angle GBs are identified in this figure. Figure 53 Top-surface GB distributions (EBSD) of ufg Cu (a) as-prepared, (b) after 400°C heat- treatment, and nt Cu (c) as-prepared, (d) after 400°C heat-treatment. Red lines represent 3 GBs; green lines represent low-angle GBs (θ = 5 – 15°), and black lines represent high angle GBs (but not 3, θ = 15 – 63°). 77 While Fig.53 provides a top-surface view of GBs, the cross-sectional view of GBs can be observed in the FIB micrographs (Fig.48, Fig.52). Table 6 presents a quantitative comparison of changes in the length fractions of 3 and coherent 3 GBs. For ufg Cu, there is a continual increase of both 3 and coherent 3 with increasing temperature, which indicates that while annealing twins were formed at elevated temperature, the fraction of ordinary (high-angle and high-energy) GBs was gradually decreased. For nt Cu, both 3 and coherent 3 were stable up to 300°C, but at 400°C, very large grains consumed some of the nt columnar grains as can be seen in Fig.52 (c) and (d). Table 6 Length fractions (EBSD) of general 3 and coherent 3 GBs on the top surfaces of ufg Cu and nt Cu before and after each heat-treatment. GB fractions of ufg Cu GB fractions of nt Cu General Σ3 Coherent Σ3 General Σ3 Coherent Σ3 As-prepared 0.35 0.07 0.50 0.38 200 C 0.51 0.21 0.50 0.38 300 C 0.66 0.33 0.51 0.39 400 C 0.71 0.39 0.42 0.31 The GB migration rate of metallic films at elevated temperature mainly depend on the minimization of GB energy, surface energy, and strain energy [53, 54]. GB energy minimization should be the major driving force for the grain growth of as-prepared ufg Cu, given that it had a large amount of ordinary GBs with few low-energy CTB. Table 6 shows that the coherent 3 (or CTB) fraction in ufg Cu increased from 0.07 to 0.21 at 200°C and had a continual increase up to 0.39 after 400°C heat-treatment, which 78 indicates the formation of annealing twins and elimination of ordinary GBs as observed Fig.53 (b). The low-mobility annealing twin boundaries formed at 200°C could stabilize big grains from further grain growth at 300 and 400°C [Fig.49 (b), (c), and (d)], as it is reported that low mobility GBs can stop grain growth by the pinning effect [7, 130-134]. In contrast, the initial microstructure of nt Cu had a high 3 GB fraction (0.50), and more importantly the coherent 3 fraction (0.38) was significantly higher than that of ufg Cu (0.07). These coherent 3 GBs on the top surface of nt Cu can be attributed to the presence of highly nt columnar grains in the cross-section as shown in Fig.48 (d). Highly nt columnar grains enhances the thermal stability of nt Cu in two ways: First, due to the compact atomic arrangement and low energy of coherent nt GBs, they are much less mobile compared with ordinary GBs at elevated temperature, and thus contribute to the thermal stability at 200 and 300°C [27, 130, 135]. Second, nt GBs intersect with columnar GBs and form many triple junctions as shown in Fig.48 (d). Saldana et al. [11] reported that the triple junctions formed by CTBs intersecting with other GBs could constrain GB migration, and thus help to obstruct high-temperature grain growth. After 400°C heat-treatment, very big grains consumed some of the nt columnar grains in nt Cu. Some GBs in nt Cu became mobile at 400°C [27], and consumed both the nanotwins and columnar GBs, thus the total GB energy of nt Cu was further reduced. The possible reasons for this rapid GB migration regarding to grain orientation will be discussed in the next section. 79 4.5 Grain orientation Figure 54 XRD peaks of (a) ufg Cu and (b) nt Cu before and after each heat-treatment. “R.T.” refers to room temperature. Fig.54 (a) and (b) show the XRD peaks of ufg and nt Cu before and after each heat-treatment. It should be noted that XRD surveyed relatively large surface areas (> 1 cm 2 ) on the samples compared with EBSD (typical scan area < 1 mm 2 ), but it can only detect a limited amount of grain orientations due to the X-ray peak formation mechanism discussed in Section 3.1.1. Fig.54 (a) illustrates that the grain orientation of ufg Cu had a drastic change to {100} after 200 C heat-treatment, and higher-temperature (300, 400 C) heat-treatment strengthened the {100} grain orientation of ufg Cu. On the other hand, for nt Cu, its top-surface grains are strongly {111} textured up to 300 C, but the {111} texture was largely weakened after 400 C heat-treatment. This XRD observation in grain 80 orientation changes is consistent with the previous result which demonstrated that ufg Cu and nt Cu lost their thermal stability at 200 and 400 C, respectively. In addition to XRD, EBSD was applied to thoroughly characterize the grain orientations before and after each heat-treatment. Fig.55 illustrates the top-surface EBSD grain orientation maps at various temperatures for ufg Cu and nt Cu foils; each color represents a specific grain orientation as illustrated in the triangle at the bottom right of the figure, and the {h k l} orientation means that the crystal direction <h k l> is normal to the foil surface. Fig.55 (a1) shows that the as-prepared ufg Cu had nearly random grain orientations. However, the random grain orientations changed and concentrated around {100} and its annealing twin {221} after 200°C heat-treatment as observed in Fig.55 (a2). This drastic change in grain orientation could be partly due to the fact that randomly oriented grains are reported [8, 136] to have generally more disordered and higher energy GBs than textured grains, and therefore could be less thermally stable. Upon 300 and 400°C heat-treatments [Fig. 55 (a3) and (a4)], {100} and its annealing twin {221} were still present and became more concentrated, which indicates that a bigger percentage of grains oriented to these two twinned orientations. To quantify the grain orientation evolution, Table 7 presents the area fractions of {100}, {111}, and {221} grain orientations before and after each heat-treatment. We can see that for ufg Cu, the {100} and {221} grain fractions continuously increase with increasing temperature, which is consistent with the results in Fig.55 (a2), (a3), and (a4). 81 Figure 55 Top-surface EBSD inverse pole figure maps of ufg Cu (a1) as-prepared, and after (a2) 200°C, (a3) 300°C, (a4) 400°C heat-treatments; nt Cu (b1) as-prepared, and after (b2) 200°C, (b3) 300°C, (a4) 400°C heat-treatments. The inset of each map is the inverse pole triangle, and the triangle at the bottom right indicates the crystallographic orientation of the maps and insets. 82 Table 7 Area fractions (EBSD) of major orientations on the top surfaces of ufg Cu and nt Cu before and after each heat-treatment (orientations deviated less than 7° from the exact major orientation are counted as this major orientation). Orientation fractions of ufg Cu Orientation fractions of nt Cu {100} {111} {221} {100} {111} {221} As-prepared 0.020 0.092 0.045 0.002 0.779 0.025 200C 0.328 0.022 0.184 0.002 0.778 0.025 300C 0.350 0.014 0.197 0.001 0.772 0.029 400C 0.421 0.008 0.245 0.260 0.538 0.043 For nt Cu, one must study two distinct stages: below and above 400°C. Below 400°C, the {111} nt columnar structure was stable as shown in Fig.55 (b1), (b2), (b3) and Table 7. This {111} texture contributes to the thermal stability in two ways: first, it is reported [8, 136] that textured microstructure generally has GBs with lower energy, and thus higher thermal stability. In our study, the as-prepared textured nt Cu had a higher fraction of low-angle GBs than as-prepared randomly oriented ufg Cu, which could lead to lower overall GB energy and thus less driving force for grain growth. Second, {111} oriented grains have minimum surface energy, and so it would be thermally stable considering surface energy minimization [55, 56]. However, this should be a minor factor since surface energy minimization is only dominant in thin (t < 5 m) metallic films, but our nt Cu sample is much thicker (t  65 m) [53, 54]. Above 400°C, Table 7 indicates that while {111} grain fraction decreases from 0.772 to 0.538, there is a big increase in the {100} grain fraction from 0.001 to 0.260. This rapid {100} grain growth [also see Fig.55 (b4)] seems to be facilitated by the enhanced GB mobility at 400°C [27], to the extent that at 600°C most {111} nt columnar grains were consumed. 83 Fig.56 (a) and (b) illustrate the grain morphology and grain orientation on the top surface of nt Cu after 600°C heat-treatment, and it can be seen that there were few {111} columnar grains left and {100} grain orientation prevailed over {111} after heat- treatment. Part of the reason for this rapid {100} grain growth in nt Cu above 400°C could be due to the highly mobile GBs of {100} grains. The GBs separating {100} and {111} grains were mostly high angle (θ > 50°), highly mobile GBs, which facilitated the rapid {100} grain growth during heat-treatments above 400°C. Figure 56 Top-surface (a) FIB micrograph and (b) XRD peaks of nt Cu after heat-treatment at 600°C. An interesting point to notice is that {100} grain orientation was favored in both ufg Cu and nt Cu foils after 400°C heat-treatment as indicated in Table 7. Previous study [58] on fcc films attached to a rigid substrate provided calculation which indicates that {100} grain has a minimum strain energy under biaxial strain during heating and would 84 grow favorably to consume grains of other orientations in order to minimize the elastic strain energy. However, given that the Cu foils in this study are free-standing, it is not clear if this strain energy minimization argument still holds during heat-treatment. Furthermore, the instability of nt Cu at 400°C is in contrast to the findings of Anderoglu et al. [8] who showed higher thermal stability in their nt Cu up to 800°C. Nevertheless, their nt Cu had a low-angle (  9°) columnar GB, while our nt Cu had many ordinary high-angle (typically 35 – 55°) columnar GBs which could provide higher driving force for GB migration. It is also suggested that by decreasing twin thickness, the number of triple junctions increases, so that the thermal stability of nt Cu could be improved. But the total GB energy increases with decreasing twin thickness, which can reduce thermal stability. So there could be a certain twin thickness that allow for optimum thermal stability. 4.6 Conclusions In this study, the effect of nanotwins on thermal stability was evaluated. Two distinct ufg Cu samples (with and without nanotwins) processed by magnetron sputtering were compared, and it was observed that the presence of nanotwins appears to change the initial GB structure and grain orientation. The sample with a highly nt structure had a higher number of CTBs and a strong {111} texture, while the other sample had few CTBs and a more randomly distributed orientation. The higher thermal stability of nt Cu was attributed to the presence of low-energy CTBs, large number of triple junctions, and 85 strongly textured {111} grain structure. However, heat treatment at 400 °C provided sufficient GB mobility for rapid {100} grain growth. Therefore, it is concluded that although the highly nt structure is only stable up to 300 °C, the presence of nanotwins does improve the thermal stability by modifying the initial GB structure and grain orientation. 86 Chapter 5 The effect of nanotwins on the corrosion behavior of Cu As previously stated in Section 2.5.2, the role of nanotwins in corrosion behavior is not well understood. In this work, four types of Cu samples containing different fractions of nt grains were characterized and corroded in order to explore how the presence of nanotwins could affect the microstructure and subsequent corrosion behavior of Cu. The major part of this corrosion study was published in Acta Materialia under the title “The effect of nanotwins on the corrosion behavior of copper” by Y. Zhao, I.C. Cheng, M.E. Kassner, and A.M. Hodge [137]. 5.1 Experimental procedures High purity Cu (99.999%) foils (  25 µm thickness) having different fractions of nt columnar grains were synthesized by magnetron sputtering following procedures described elsewhere [125, 128]. In addition, mc Cu samples were prepared by heating the as-sputtered Cu (without nanotwins) at 600 C for 3 hours in a MTI-GSL1100X furnace under vacuum (  3 × 10 -5 Torr). Each sample was partially mounted in epoxy as described in Section 3.2.2, and the exposed areas of the samples for polarization and immersion corrosion tests were  0.36 cm 2 and  0.64 cm 2 , respectively. All samples were ultrasonically cleaned in acetone and ultra-pure water before the corrosion tests, and rinsed gently with water after corrosion. The corrosion solution is artificial seawater: 3.5% (35 g/L) NaCl solution with 8.0  0.1 pH value. 87 Polarization tests were performed in a three-electrode electrochemical cell. The (1) working electrode was the Cu sample. The (2) reference electrode was silver/silver chloride (Ag/AgCl) which was placed close to the working electrode. The (3) counter electrode was a platinum wire. These three electrodes were connected to a potentiostat (Gamry Reference 3000) which performed two types of polarization tests: linear polarization and potentiodynamic polarization. More details about the set-up and basics of electrochemical plorization were described in Section 3.2.2. In linear polarization tests, the corrosion potential (U corr ) of each Cu sample was stabilized in the solution for 15 minutes before polarization in order to achieve a linear relationship around U corr . The potential was then swept from – 10 mV to 10 mV (vs U corr ) at a rate of 0.1 mV/s. For potentiodynamic polarization tests, the potential was polarized from – 450 mV to 450 mV (vs Ag/AgCl) with a scan rate of 0.5 mV/s immediately after the sample was placed in the solution. Immediate polarization of the potential minimizes the reaction between the Cu sample and corrosion solution, so that the cathodic peak [138-140] which perturbed the Tafel linear region could be minimized. Furthermore, Cu samples after linear polarization were subsequently polarized from – 450 mV to the passivation point (see Fig.44) at a scan rate of 0.5 mV/s in order to observe the passive layers. The curves of linear and potentiodynamic polarization tests were analyzed by Gamry Echem Analyst software. At least two samples from each type (columnar highly nt, columnar partially nt, columnar non-nt, and non-columnar mc) of Cu foil were tested in either linear polarization or potentiodynamic polarization. The polarization resistances (R p ), Tafel slopes (β c , β a ), and passivation current densities (I pass ) were measured. The 88 corrosion current densities (I corr ) of each type of Cu were calculated through the Stern- Geary equation by inputting the measured β c , β a and each R p . The standard deviations of R p , I corr and I pass were calculated. For immersion corrosion tests, two samples from each type of Cu foil were placed in 3.5% NaCl solution containers with the lid closed for a total of 10 days. The samples were ultrasonically cleaned, and their passive layers were evaluated after immersion corrosion. Since the passive layers of the two samples from each type of Cu foil were similar, only one of them is presented. The cross-sectional and top-surface (corrosion surface) microstructures of the as- prepared Cu samples were characterized by FIB and EBSD, respectively. The details of EBSD analysis are consistent with that described in Section 4.1. A {hkl} grain orientation means that its crystal direction <hkl> is normal to the top surface, and crystal orientations within 7  of a specific {hkl} are considered belonging to that {hkl} grain orientation. After corrosion, the top-surface and cross-sectional view of passivated samples during polarization tests were characterized by SEM and FIB, respectively. The passive layers formed during immersion corrosion were also characterized by SEM. EDS was used to obtain the compositions at the microscopic (10 – 100 µm) level, while XRD was used to detect the chemical compounds and crystal orientations of corrosion products at the macroscopic (  10 mm) level. 89 5.2 As-prepared samples Figure 57 FIB cross-sectional micrographs of as-prepared (a) highly nt, (b) partially nt, (c) non- nt, and (d) mc Cu samples. The arrows indicate growth direction and point to the top surfaces which will be exposed to the corrosion solution. The cross-sectional FIB micrographs of the as-prepared Cu samples are illustrated in Fig.57. The arrow in each figure indicates the growth direction of Cu foil during magnetron sputtering, and it points to the top surface which will be exposed to the corrosion solution. Fig.57 (a) shows the Cu sample with nanotwins within over 90% of the columnar grains, and is denoted as “highly nt Cu”. Note that some columnar grains in Fig.57 (a) do not show nanotwins due to their unfavorable orientations for ion channeling contrast, and the FIB ion channeling mechanism is discussed in Section 3.1.4. Fig.57 (b) shows nanotwins in  60% of columnar grains, and this is denoted as “partially nt Cu”. Fig.57 (c) shows Cu with ~ 5% nt columnar grains, and is denoted as “non-nt Cu”. These percentages represent the number fractions of nt columnar grains, which were obtained by surveying around 100 columnar grains at different locations for each type of Cu 90 sample. The columnar width of highly nt, partially nt, and non-nt Cu were measured to be  0.40,  0.34, and  0.30 µm, respectively. Since these three types of Cu samples have columnar grains, they are altogether called “columnar Cu”. The columnar widths of these Cu samples are similar and within the ufg range (0.1 µm < grain size < 1 µm). Fig.57 (d) illustrates that mc Cu has micron-sized grains and the grain shape is irregular instead of columnar. In addition, the EBSD top-surface grain orientation maps are presented in Fig.58. The triangle legend at the bottom right of Fig.58 indicates the crystallographic orientations in each map and inset, and the black, green, and red GBs represent high- angle, low-angle, and general twin boundaries (including both CTBs and incoherent twin boundaries), respectively. Notice that the micron bar of Fig.58 (d) is different from those of Fig.58 (a), (b), (c) due to the much bigger grain size of mc Cu. The top-surface grain sizes of these samples are listed in Table 8, which shows that columnar Cu samples have grain sizes of around 0.4 µm, but mc Cu has a larger grain size of 8.6 µm. Table 8 also lists the quantitative information of grain orientations and GBs obtained from the EBSD maps shown in Fig.58. It demonstrates that highly nt Cu has the highest fraction of {111} grain orientation and CTBs on the top surface. This is largely due to the presence of nanotwins. Since the orientation plane of nanotwin GB is {111}, and given [Fig.57 (a)] that these planes are parallel with the top surface, each nt columnar grain should have a {111} grain orientation on the top surface. Consequently, highly nt Cu which mostly consists of nt columnar grains has a {111} texture. 91 Figure 58 EBSD top-surface grain orientation map overlapped with GBs: as-prepared (a) highly nt, (b) partially nt, (c) non-nt, (d) mc Cu. The inset in each map is the inverse pole triangle, and the triangle legend at bottom right indicates specific crystallographic orientations in the maps and insets. The legend at top right correlates twin boundaries, high-angle and low-angle GBs with red, black, and green lines, respectively. In a, b and c, the observed twin boundaries are mostly due to the nanotwins intersecting with the surface; the curled shape of these twin boundaries is due to the surface roughness. Furthermore, since some nanotwins in the nt columnar grains intersect with the top surface, increasing the overall numbers of nt columnar grains will increase the number of CTBs on the top surface. Compared with highly nt Cu, partially nt and non-nt Cu have gradually less nt columnar grains, so their fractions of {111} grain orientation and CTBs continually decrease as presented in Table 8. For mc Cu, Fig.58 (d) shows that its major orientations are {100} and its twinned grain orientation {221}. The mc Cu has 92 a large fraction (  0.36) of CTBs due to the formation of annealing twins during heat- treatment. Over 60% of GBs on its top-surface are ordinary GBs which are neither CTB nor low angle (< 15 ) GBs. Table 8 EBSD quantitative information of grain size, grain boundaries, and grain orientations on the top surface of the Cu samples. Grain size ( m) a Area fraction of grain orientations Length fraction of special GBs over all GBs b Length fraction of ordinary GBs over all GBs CTBs Low- angle GBs {100} {111} {221} columnar highly nt 0.52 0 0.88 0.01 0.52 0.09 0.39 partially nt 0.41 0.03 0.33 0.10 0.26 0.07 0.67 non-nt 0.34 0.02 0.08 0.04 0.07 0.05 0.88 mc 8.6 0.23 0.02 0.32 0.36 0.02 0.62 a Twin boundaries are not included in grain size measurement; b Special GBs are relatively corrosion resistant. 5.3 Polarization 5.3.1 Polarization curve analysis The polarization resistances (R p ) of Cu samples obtained from linear polarization are listed in Table 9: the R p of partially nt, non-nt, and mc Cu are similar and close to  10 kΩ∙cm 2 , while the R p (  29 kΩ∙cm 2 ) of highly nt Cu is about two times higher. 93 Table 9 Summary of polarization tests results. Polarization resistance (R p , kΩ·cm 2 ) Cathodic slope (β c , mV/decade) Anodic slope (β a , mV/decade) Corrosion current density (I corr , A/cm 2 ) Passivation current density (I pass , mA/cm 2 ) columnar highly nt 28.6 ± 7.1 194 23 0.34 ± 0.09 1.15 ± 0.20 partially nt 10.2 ± 4.4 323 23 1.08 ± 0.39 1.14 ± 0.23 non-nt 10.4 ± 0.5 206 35 1.26 ± 0.06 1.08 ± 0.27 mc 9.2 ± 2.1 172 50 1.92 ± 0.42 1.27 ± 0.18 Fig.59 shows the typical potentiodynamic curves of the four types of Cu. The major reaction of the cathodic Tafel regions is [140, 141]: O 2 + 4e - + 2H 2 O  4OH - , while the major reaction of the anodic Tafel regions is [140, 141]: Cu  Cu + + e - . Please refer to the Appendix B for the further discussions on the anodic Tafel regions. The cathodic slopes (β c ) and anodic slopes (β a ) of the linear Tafel regions were measured and listed in Table 9. Given that β c , β a , and R p are known, the corrosion current density (I corr ) can be calculated through Stern-Geary equation: I corr = β c ∙β a /[2.3∙R p ∙(β c +β a )]. Table 9 lists the values of I corr , and Fig.60 (a) and (b) plot the I corr and passivation current density (I pass ) of Cu samples during polarization, respectively. It can be seen from Fig.60 (a) that the columnar Cu samples have smaller I corr than mc Cu, and highly nt Cu has the smallest I corr among the columnar Cu samples. 94 Figure 59 Potentiodynamic polarization curves of highly nt, partially nt, non-nt, and mc Cu in the 3.5% NaCl (pH ~ 8.0) solution. The two y-axes are potentials versus Ag/AgCl and SHE reference potentials. It is interesting to note that columnar Cu samples have lower corrosion rates compared to mc Cu even though they have a much larger area of GBs which are generally susceptible to corrosion. This phenomenon, which was also reported in other Cu corrosion studies [73, 74, 76], could be primarily due to a grain size effect. Miyamoto et al. [76] explained that with smaller grain size, the Cu ions that dissolved from GBs could easily flow upon the surface of grain interior and interrupt the supply and reduction of O 2 , which results in the lower corrosion rate of smaller grain sized Cu. 95 Figure 60 (a) corrosion current density (I corr ) and (b) passivation current density (I pass ) of highly nt, partially nt, non-nt, and mc Cu samples in the 3.5% NaCl (pH ~ 8.0) solution. In addition, among columnar Cu samples, highly nt Cu has a much lower corrosion rate compared with both partially nt and non-nt Cu. This could be due to its unique GB network and the {111} texture. It is generally accepted that CTBs [19] and low-angle GBs [142] have good corrosion resistance, and thus are considered as “special 96 GBs”. From the perspective of GBE, increasing the fraction of special GBs helps to interrupt the continuity of ordinary GBs, so that the intergranular corrosion through ordinary GBs can be arrested by special GBs [19, 33, 34]. Wang et al. [50] examined the nanotwins and columnar GBs in a highly nt Cu by using TEM orientation mapping. They reported that in highly nt Cu, some of the columnar GBs [see Fig.13 (a)] have alternating segments of high-angle and low-angle GBs due to the alternating nanotwins, while some other columnar GBs [see Fig.13 (b)] are primarily high-angle GBs. These two configurations of columnar GBs in highly nt Cu are schematically illustrated in Fig.61 (a) and (b), respectively. It can be observed from the unique GB network of Fig.61 (a) that even though corrosion can initially proceed through a high-angle segment of a columnar GB, it will encounter a low-angle GB and a nt GB (which is a CTB). Since both low- angle GB and CTB have excellent corrosion resistance, the intergranular corrosion will likely be arrested [19, 34]. Thus, some columnar GBs in highly nt Cu could have improved corrosion resistance due to this unique GB network. Besides GB network, corrosion resistance is also related to grain orientation. Different grain orientations have distinct atomic binding energies on the orientation plane, so their resistance to metal dissolution in corrosive environments are also different [21, 143, 144]. The {111} orientation of fcc metals is reported [21, 143] to have higher corrosion resistance than other grain orientations, as it is the close-packed plane which has the highest atomic binding energy. Consequently, the {111} texture [as illustrated in Fig.58 (a)] of highly nt Cu can also contributes to its higher corrosion resistance. 97 Figure 61 Schematic illustration of two possible columnar grain boundary configurations within a nt Cu sample: (a) alternating high-angle and low-angle boundary segments, and (b) high-angle boundary only. The legend at bottom correlates red, black, and green lines with nanotwin, high- angle, and low-angle GBs, respectively. Different crystal orientations are illustrated by different gray-scale colors. The two arrows point along the corrosion surface. Upon further increase in potential, the current density decreases due to passivation as shown in Fig.59. The current density then reaches a minimum value called the passivation current density (I pass ) where a metal is passivated, and the values of I pass are listed in Table 9 and plotted in Fig.60 (b). Fig.60 (b) illustrates that the I pass of columnar Cu samples are similar, while the I pass of mc Cu is slightly higher, which may suggest better passivation behaviors of columnar Cu samples compared with mc Cu. Microstructural analysis of the passive layers formed on these passivated Cu will be presented in the following section. 98 5.3.2 Passive layer morphology Fig.62 shows the SEM top-surface and FIB cross-sectional micrographs of passivated Cu samples which reached the passivation current densities. Figure 62 SEM top-surface morphologies and FIB cross-sectional micrographs of passivated (a1, a2) highly nt, (b1, b2) partially nt, (c1, c2) non-nt, and (d1, d2) mc Cu during polarization in the 3.5% NaCl (pH ~ 8.0) solution. Markers “o” and “x” indicate nt and non-nt columnar grains, respectively; each arrow marker points to a specific GB. The black layer on the top of each FIB micrograph is the FIB deposited carbon layer. The actual length in the vertical direction of FIB micrographs is  1.27 times the imaged length due to the 52  tilt. 99 From Fig.62 (a1), (b1), and (c1), the passive layers formed on columnar Cu samples are generally smooth and uniform with the presence of shallow grooves. Additional corrosion products were formed on the passive layers, but no corrosion pits can be observed. In contrast, the surface of passivated mc Cu has large corrosion pits as shown in Fig.62 (d1). For the cross-sections, Fig.62 (a2), (b2), and (c2) illustrate that columnar Cu samples are covered by continuous passive layers that adhere to the Cu base metals and protect them from further corrosion. However, no continuous passive layer can be observed on the passivated mc Cu, as shown in Fig.62 (d2). Three corrosion pits of mc Cu are marked by numbers 1 - 3 in Fig.62 (d2), and it can be observed that pit #2 is not covered by corrosion products. This suggests that the corrosion products of mc Cu may not be adhesive to the base metal during polarization, and without the protection of corrosion products, the corrosion pits of mc Cu could have uninterrupted growth. These microstructural observations indicate that the passive layers of columnar Cu are more protective than that of mc Cu, which is also consistent with the slightly lower passivation current densities of columnar Cu samples shown in Fig.60 (b). The enhanced passivation property of columnar Cu could be due to their smaller grain sizes. Studies [73, 78] have shown that small grain sizes could increase the activity of surface electrons and the absorption of Cl - ions on the surface, which facilitates the formation of a continuous and protective passive layer. Moreover, Tao and Li [73] reported that small grain size could improve the adhesion of passive layers on the base metal due to the increased activity of electrons at GBs. Lu et al. [78] observed no corrosion product layer on cg Cu-20Zr after corrosion, while a continuous passive layer was formed on nc Cu- 100 20Zr under the same corrosion condition. These findings are consistent with the results of this study. 5.3.3 Corrosion paths For columnar Cu samples, Fig.62 (a2), (b2), and (c2) illustrate that the corroded columnar grains generally have a shape of hills and valleys. The “valleys” are always connected to the columnar GBs, which suggests that corrosion proceeded through columnar GBs. It is evident especially for highly nt Cu [Fig.62 (a2)] that some columnar GBs (indicated by the arrows) were preferentially corroded. As previously discussed in Section 5.3.1, some columnar GBs in highly nt Cu have alternating segments of high- angle and low-angle GBs, which should provide relatively high intergranular corrosion resistance. However, as illustrated in Fig.61 (b), some other columnar GBs in highly nt Cu are primarily high-angle GBs, which may be vulnerable to corrosion and can be selectively attacked. Therefore, in order to achieve better intergranular corrosion resistance, the high-angle columnar GBs in highly nt Cu need to be modified by low- angle GB segments. Furthermore, in Fig.62 (a2) and (b2), some nt and non-nt columnar grains are marked by “o” and “x”, respectively. It can be observed that these non-nt columnar grains were preferentially corroded compared with their neighboring nt columnar grains, which suggests that these two types of columnar grains could form a localized galvanic 101 couple where nt columnar grain serves as the noble side. These observations demonstrate the improved corrosion resistance of nt columnar grains. For mc Cu, Fig.62 (d2) indicates that it has more substantial corrosion pits compared with columnar Cu samples [Fig.62 (a2), (b2), and (c2)]. Moreover, pit #1 and pit #3 in Fig.62 (d2) possibly proceeded through the two GBs indicated by the arrows, as their lowest points are connected with these two GBs. These observations are consistent with the studies [74, 76] which suggest that the Cu sample with larger grain size is more likely to have preferential corrosion at GBs and subsequent pitting due to its lower area fraction of GBs over grain interior. 5.4 Immersion corrosion In addition to polarization, immersion corrosion tests were carried out using a 3.5% NaCl (pH  8) solution for 10 days. Fig.63 shows the top-surface topographies of these samples after immersion corrosion, and additional corrosion products can be observed on top of the passive layers. Since these additional corrosion products block the view of the passive layers, it is difficult to evaluate the intactness of the passive layers. Therefore, these Cu samples were ultrasonically cleaned in order to remove these additional corrosion products. 102 Figure 63 SEM top-surface topographies of the passive layers on (a) highly nt, (b) partially nt, (c) non-nt, and (d) mc Cu samples after immersion corrosion in 3.5% NaCl (pH ~ 8.0) solution. Fig.64 shows the passive layers after ultrasonic cleaning. Highly nt Cu [Fig.64 (a)] has a relatively intact and uniform passive layer, and no large corrosion pits were observed. In contrast, large corrosion pits were observed on the passive layers of partially nt and non-nt Cu as shown in Fig.64 (b) and (c), respectively. Thus, although the passive layers [Fig.62 (a1), (b1), and (c1)] of the columnar Cu samples formed during polarization are similar, highly nt Cu has a less defective passive layer than the other columnar Cu samples after immersion corrosion. For mc Cu [Fig.64 (d)], its passive layer contains large corrosion pits, which is comparable to the pits formed on passivated mc Cu [Fig.62 (d1)] during polarization. 103 Figure 64 SEM top-surface topographies of the passive layers on (a) highly nt, (b) partially nt, (c) non-nt, and (d) mc Cu samples after immersion corrosion in 3.5% NaCl (pH ~ 8.0) solution and subsequent ultrasonic cleaning. The high pitting resistance of highly nt Cu is consistent with other corrosion studies on Ni [15-17] and Al [18] with nanotwins. Meng et al. [15] and Sun et al. [16] further indicated that the passive layer formed on nt structure had a very low diffusion coefficient, which was responsible for its significantly improved pitting corrosion resistance. However, it was not explicitly explained how nanotwins could help to form such a protective passive layer. The protection of the passive layer formed on highly nt Cu in this work will be further discussed in the following section. 104 5.5 Composition of passive layers Fig.65 presents the normalized XRD peaks of corroded Cu samples. Note that the XRD peaks of Cu were removed in order to facilitate the comparison of corrosion products, and the {200} peak of Cu 2 O was also removed due to its proximity to the {111} peak of Cu. As shown in Fig.65 (a), the major corrosion products of passivated columnar Cu samples during polarization are cuprous chloride (CuCl) and cuprous oxide (Cu 2 O). EDS also detected both Cl and O elements in the passive layer and so confirmed the existence of CuCl and Cu 2 O. These passive layer compositions are consistent with the literature review [141] which indicates that CuCl and soluble cuprous complex (CuCl 2 - ) can be formed via reactions: Cu + + Cl -  CuCl, Cu + +2Cl  CuCl 2 - , and Cu 2 O could be subsequently formed by the precipitation of CuCl 2 - : 2CuCl 2 - + 2OH -  Cu 2 O + H 2 O + 4Cl - . For passivated mc Cu, the XRD peaks of CuCl and Cu 2 O are much lower than those of passivated columnar Cu samples, as no observable passive layer was formed on mc Cu as shown in Fig.62 (d2). In comparison with polarization, the passive layer composition of immersion corroded Cu is different. The immersion corrosion products were Cu 2 O, copper chloride hydroxide [Cu 2 (OH) 3 Cl], and no CuCl could be detected, which is also consistent with the immersion studies of Cu under NaCl [145] and seawater [146, 147] environments. After ultrasonic cleaning, most Cu 2 (OH) 3 Cl corrosion products were detached from the passive layers, and only Cu 2 O can be detected as shown in Fig.65 (b). The reason for the compositional difference between passive layers formed during polarization and immersion corrosion could be that polarization is a short-term corrosion, and factors such 105 as pH value and ion concentrations can change dynamically and may not be evenly distributed in the solution, while immersion corrosion is a long-term corrosion. Care must be exercised when using electrochemical polarization to predict the passive layer and long-term corrosion behavior. Figure 65 Normalized XRD peaks showing the corrosion products of (a) passivated Cu, and (b) immersion corroded and subsequently ultrasonic cleaned Cu samples. The solid squares and circles on top of the peaks mark CuCl and Cu 2 O, respectively; the peak orientations are also marked. 106 In spite of that, both types of corrosion methods indicate that the Cu 2 O {111} peak on highly nt Cu is much higher than those of the other samples as illustrated in Fig.65 (a) and (b). This suggests that there could be a preferential growth of {111} oriented Cu 2 O passive layer on highly nt Cu. Studies [148, 149] indicated that there was an epitaxial growth of Cu 2 O {111} on top of Cu {111}, and Kunze et al. [150] further suggested that Cu 2 O {111} was stable due to the high cohesive energy of its equivalent monolayers in the {111} orientation. Therefore, the stable and protective Cu 2 O {111} passive layer of highly nt Cu could be originated from its {111} texture. 5.6 Conclusions Columnar Cu samples with different fractions of nanotwins and mc Cu samples were prepared, characterized and tested by both polarization and immersion corrosion in 3.5% NaCl solution (pH  8.0). Highly nt Cu has a much lower corrosion current density compared with the other types of Cu samples, which is attributed to its nanotwin-induced GB network and the unique {111} texture. Moreover, columnar grains with nanotwins were observed to be more corrosion resistant than those without nanotwins. Although the passive layers formed during polarization and immersion corrosion tests are not identical, both tests indicate that the passive layer of highly nt Cu is relatively stable and has improved pitting resistance compared with the other Cu samples. This protective passive layer could be related to the preferential growth of Cu 2 O {111} on top of the highly textured {111} nt Cu. Overall, this study indicates that the presence of nanotwins 107 appears to modify the microstructure which subsequently improves the corrosion properties of Cu. However, further synthesis control is needed in order to minimize the vulnerability due to the presence of high-angle columnar GBs and columnar grains without nanotwins. 108 Chapter 6 The effect of grain boundary plane orientation on the sensitization of an Al-Mg alloy Al-Mg alloys having more than 3 wt% Mg content can lead to Mg preferentially diffusing to the GBs and forming an intergranular β phase (Al 3 Mg 2 ). The intergranular β phase corrodes preferentially compared with the Al-Mg matrix in most corrosive environments, which leads to intergranular corrosion and stress corrosion cracking [1, 2]. This phenomenon is also known as “sensitization” of Al-Mg alloys [2, 87]; background details were introduced in Section 2.6. Nevertheless, how GB plane orientation affects β precipitation has not been directly explored in the previous research. Therefore, this work focuses on the role of GB plane orientation in the sensitization behavior of Al 5456 for both low-angle and high-angle GBs. β precipitation of a special columnar GB structured Al 5456 with nc grains was also examined. The major part of this sensitization study was submitted to Scripta Materialia under the title “The role of grain boundary plane orientation in the β phase precipitation of an Al-Mg alloy” by Y. Zhao, M.N. Polyakov, M. Mecklenburg, M.E. Kassner, and A.M. Hodge. 6.1 Experimental procedures As-received Al 5456-H116 (denoted as “AR-5456”) was purchased from Aleris Aluminum Koblenz GmbH, and H116 refers to a strain-hardening for obtaining desired strength [87]. Magnetron sputtering targets were machined from the AR-5456 to synthesize the as-sputtered nc Al 5456 (denoted as “AS-5456”) with a ~ 20 m foil 109 thickness. Both AR-5456 and AS-5456 were heated at 175°C for 7 days under ~ 2 × 10 -6 Torr vacuum in order to evaluate their sensitization behaviors. The grain size of AR- 5456 was determined by EBSD, while that of AS-5456 was obtained by TEM and transmission-EBSD (details in Section 3.1.3). After heat-treatment, the top surfaces of the samples were etched using 10% phosphoric acid (H 3 PO 4 ) at 35°C for 1 minute so that the β phase could be revealed [100, 124]. This etching condition not only effectively revealed the β phase, but also had minor etching effect on the areas without β phase, and thus can exclusively reveal the sensitized GBs as presented in Section 3.2.4. SEM was used to observe the etched GBs. Lamella containing cross-sections of the etched areas in AR-5456 and AS-5456 were prepared using standard FIB procedures, then imaged and analyzed using TEM techniques. The detailed procedures used by Unocic et al. [124] are shown in Fig.66: first, the GBs were etched by phosphoric acid in Fig.66 (a); second, Fig.66 (b) illustrates that a specific etched GB was selected and thinned into a foil by FIB. Note that the selected GB traverses the foil thickness and is perpendicular to the horizontal axis of the foil; third, the foil lifted-out by FIB was then placed in a TEM to observe the etched GB and its β precipitation in Fig.66 (c). This work followed similar procedures to observe the intergranular β phase, and the existence of β phase in FIB specimens was confirmed by EELS and EDS acquired using STEM methods. 110 Figure 66 Method for the observation of intergranular β phase: top-surface (a) phosphoric etching of the β phase, (b) FIB milling and thinning, and cross-sectional (c) TEM observation of the FIB lift-out foil. Adapted from [124]. 6.2 As-prepared samples Fig.67 (a) shows the top-surface EBSD map of AR-5456, and it can be observed that AR-5456 has preferred {110}, {431}, and {843} grain orientations. The cross- sectional TEM dark-field images [Fig.67 (b)] illustrate that AS-5456 has columnar grains with much smaller grain size compared with AR-5456. 111 The grain sizes and elemental compositions of AR-5456 and AS-5456 before and after heat-treatment are listed in Table 10. The AR-5456 sample had an average grain size of  21 m with ~ 5.2 wt% Mg, while AS-5456 had a much smaller average grain size of  0.13 m and the Mg content decreased to ~ 4.3 wt%. The Mg loss in AS-5456 was possibly due to the high vapor pressure of Mg during magnetron sputtering. Table 10 also indicates that there was no substantial change in the grain size and Mg content of both samples after heat-treatment. Figure 67 (a) Top-surface EBSD map of AR-5456, and (b) cross-sectional dark-field TEM graphs of AS-5456. The triangle legend in (a) correlates the color of the EBSD map with grain orientations. 112 Table 10 Grain sizes and elemental compositions of AR-5456 and AS-5456 before and after heat-treatment. Al-Mg samples Heat-treatment (175°C × 7 days) Grain size ( m) Elemental composition (wt%) Mg Al Other elements a AR-5456 before 20.9 5.2 93.4 balance after 21.8 5.1 93.2 AS-5456 before 0.13 4.3 94.2 after 0.14 4.4 92.6 a Other elements include: Mn, Fe, Si, Cr, Zn, Ti, Cu, etc. 6.3 Sensitization behavior of as-received (AR) 5456 Figs.68 (a) and (b) show the SEM topography and EBSD grain orientation map of the identical etched area of the heat-treated AR-5456. We define the “Z” axis as the axis perpendicular to the top-surface as shown by the legend in Fig.68. In (a), the etched GBs are highlighted, which signifies the presence of intergranular β phase; (b) correlates the etched GBs in (a) noting their specific GB misorientations. The red, yellow and black lines in Fig.68 (b) represent low-CSL (Σ≤29), low-angle (θ≤15°), and high-angle (θ>15°) GBs, respectively. The sample had ~ 3% low-CSL (Σ≤29) GBs, and many of them were etched. This indicates that low-CSL GBs do not necessarily possess resistance to β precipitation, which is in contrast to the suggestion of Kaigorodova [23]. Figs.68 (a) and (b) also illustrate four particular GBs (GB-1, GB-2, GB-3, GB-4), and the neighboring grains of GB-1 and GB-2 are marked by the circled numbers. The misorientations and etching behaviors of the four GBs are listed in Table 11. It can be observed that low- angle GB-2 and GB-4 were etched, suggesting that low-angle GBs can have β precipitation. This is in agreement with Scotto D’Antuono et al. [24], but in contrast to 113 Davenport et al. [22] who suggested that low-angle GBs are resistant to β precipitation. For high-angle GBs, Fig.68 (a) shows that although GB-1 was etched, GB-3 was not. This suggests that some high-angle GBs are immune to β precipitation, which was also observed by Davenport et al. [22] and Kaigorodova [23]. Davenport et al. [22] further reported that GBs with similar misorientations could have different tendencies for etching. Therefore, in addition to GB misorientation, GB plane orientation may also play a role in β precipitation. This is an emphasis for the present study. Figure 68 Top-surface (a) SEM topography and (b) EBSD grain orientation map of the identical area of the heat-treated/etched AR-5456. The white rectangles in (a) are the FIB lift-out areas, and the circled numbers mark four distinct grains. Red, yellow and black lines in (b) represent low-CSL, low-angle and high-angle GBs, respectively. The bottom-right triangle correlates each color in the EBSD map with a specific grain orientation. 114 Table 11 Different GB misorientations and the corresponding etching behaviors. Misorientation angle θ (°) Misorientation axis Etching behavior GB-1 48.7 <17 8 3> etched GB-2 8.1 <7 6 4> etched GB-3 54.0 <7 7 1> not etched GB-4 11.5 <25 10 7> etched To explore the influence of GB plane orientation on β precipitation, FIB prepared lamella were extracted from the regions indicated by white rectangles in Fig.68 (a), each region contained a GB (indicated as GB-1 and GB-2 in the figure). TEM images of the lamella are shown in Figs.69 (a) and (b). The “Z” axis is now upward in the plane of Fig.69 due to the view rotation from Fig.68. The circled numbers correspond to the same grains as shown in Fig.68, and the region between these grains is the GB covered by β phase. The red arrows indicate the etching direction, and the etch widths (or β phase thicknesses on the surface) of GB-1 and GB-2 are ~ 600 and ~ 125 nm, respectively. Figs.69 (a) and (b) also show four Mg maps (the dashed box areas) obtained by EELS core-loss spectrum imaging; the brighter regions contain ~ 36 – 40 wt% Mg, which confirms the presence of the β phase (37.5 wt% Mg). In addition, it is observed that the β phase thickness largely varies along GB-1 and GB-2. Given that the GB misorientation is roughly constant along a single GB, the variation in β phase thickness could be due to the change in GB plane orientation. We then estimated the four GB planes with the black dashed traces (p1, p2, p3, p4) along GB-1 and GB-2 as shown in Figs.69 (a) and (b). The GB associated with each trace is 115 perpendicular to the graphic plane, which is a consequence of the orientation between the GB and the FIB prepared lamella in Fig.68. The GB plane orientations were then determined by measuring the directions perpendicular to the traces p1, p2, p3, p4 in Figs.69 (a) and (b). Detailed measurement procedures are explained in Appendix C. Table 12 lists the measured GB plane orientations which are also illustrated in the inverse pole triangle shown in Fig.69 (c): the darker dots represent thicker β phase, while the hollow dots are GB planes without β precipitation. The β phase thickness of each GB plane is also listed in Table 12. It can be observed that GBs with thicker β phase (p1and p3) have at least one plane orientation that is close to {110}, which may suggest the vulnerability of near {110}-oriented GB planes to β precipitation. Figure 69 Cross-sectional TEM micrographs of the two GBs lifted out by FIB in Fig.68: (a) GB- 1, and (b) GB-2. Mg composition maps of the dashed box areas were obtained by EELS. The circled numbers mark the four grains, and p1 – p4 indicate the estimated GB plane traces. (c) GB plane orientations of p1 – p4 in an inverse pole triangle: darker dots represent thicker β phase, and hollow dots refer to no β precipitation. 116 Table 12 β phase thicknesses on the different GB planes of the sensitized AR-5456. GB planes GB plane orientations a β phase thickness (nm) GB-1 (high angle) p1 G1: {10 5 1} G2: {7 6 1} 580 p2 G1: {12 10 9} G2: {9 3 2} 75 GB-2 (low angle) p3 G3: {21 16 1} G4: {9 8 2} 130 p4 G3: {14 7 1} G4: {7 4 1} no β phase a G1 – G4 are the four grains marked by the circled numbers in Fig.1 and Fig.2. 6.4 Sensitization behavior of as-sputtered (AS) 5456 We additionally explored the sensitization behavior of nc AS-5456 which had columnar grains and a {111} texture (not shown) on the top surface. Fig.70 shows different views including the top surface (a) and cross-sections (b – d) of the heat- treated/etched AS-5456. It can be observed from Fig.70 (a) that the smooth surface of the heat-treated/etched AS-5456 is in contrast to the clear intergranular corrosion of the heat-treated/etched AR-5456 in Fig.68 (a). This demonstrates the minimal sensitization of AS-5456. The inset of Fig.70 (a) presents a much higher magnification image which reveals small, discontinuous etch pits (average size ~ 70 nm) on the top surface. Fig.70 (b) shows a cross-sectional TEM image of columnar grains, and it can be observed that most columnar GBs are tilted α degrees (10 – 25  ) from the <111> crystallographic direction of the {111} textured AS-5456. Fig.70 (c) shows four typical columnar GBs denoted as GB-I, GB-II, GB-III, and GB-IV. Etch pits marked by the red arrows can be observed on the top surfaces of GB-III and GB-IV, and STEM-EDS detected a region indicated by the yellow arrows in GB-IV that has an enrichment of ~ 11 wt% Mg. Nevertheless, no GBs in (c) exhibit any β phase. Although AS-5456 (4.3 wt% Mg) has a 117 slightly lower Mg content than AR-5456 (5.2 wt% Mg), it can be directly compared to Al 5083 with a similar Mg content (4.0 – 4.9 wt%). In contrast to the intergranular β precipitation of Al 5083 [2, 89, 90], the AS-5456 demonstrated excellent resistance to β precipitation under similar elevated-temperature conditions. Figure 70 Top-surface (a) SEM and cross-sectional (b) TEM, (c) magnified TEM, and (d) transmission-EBSD graphs of heat-treated/etched AS-5456. In (c), the yellow arrows of GB-IV refer to one Mg-rich region. The grain orientations (normal to the graphic plane) and GB types in (d) are defined by the legends in Fig.68. In order to understand the main factor that could lead to the improved sensitization resistance of AS-5456, the columnar GB types were identified using 118 transmission-EBSD as shown in Fig.70 (d). The red, yellow, and black lines represent low-CSL, low-angle, and high-angle GBs, respectively. It can be observed that the majority of GBs are high-angle GBs. Thus, GB misorientation may not contribute to the sensitization behavior of AS-5456. From the perspective of grain size, although AS-5456 has nc grains, it may have little influence on the β precipitation. According to the diffusion model proposed by Porter and Easterling [97] (Equation 5 in Section 2.6.2), the growth rate of β phase is inversely proportional to β phase thickness at the beginning, but will have an additional drop if the Mg diffusion zone reaches the limit (typically half of the grain size). Therefore, a smaller grain size may result in a decrease in β phase growth rate. We thus estimated the maximum β phase thickness of AS-5456 before its small grain size could affect β phase growth rate, and calculated it to be  3.3 nm through the Equation 4 of Section 2.6.2. Since β phase was actually not observed, grain size may have little effect on the sensitization behavior of AS-5456. One factor which could influence β precipitation is dislocation density. The AS- 5456 was magnetron sputtered, and this process may lead to lower dislocation density compared with the conventional strain-hardened AR-5456 [128]. Lower dislocation density could then restrain the dislocation pipe diffusion of Mg towards GBs and thus decrease the overall β precipitation [89, 92]. However, even if we assumed that AS-5456 was dislocation-free, one would still expect an estimated β phase thickness to be  29 nm through the volume diffusion equation [89, 92, 97, 151]. Therefore, given that no β phase was observed, there should be other factors in addition to dislocation density that could improve the sensitization behavior of AS-5456. As discussed earlier, GB plane 119 orientation may play a significant role in β precipitation in addition to GB misorientation. In the case of AS-5456 which has columnar grains and a {111} texture, it could have less variation in GB plane orientations compared to a sample with randomly-oriented equiaxed grains. Thus, the columnar GB in AS-5456 could have plane orientations that are less vulnerable to β precipitation. However, a detailed assessment of these columnar GB orientations needs to be performed in order to understand the effect of GB plane orientation on the sensitization behavior of AS-5456. 6.5 Conclusions In summary, the variation of intergranular β phase thickness in the sensitized AR- 5456 demonstrates the dependence of β precipitation on GB plane orientation. Specifically, GBs that have plane orientations near to {110} may facilitate β precipitation. Moreover, AS-5456 possessed excellent sensitization resistance, which may be due to its columnar GB plane orientations. More research is needed to identify the columnar GB plane orientations of AS-5456 in order to further explore how intergranular β precipitation could be diminished by altering the GB plane orientation. 120 Chapter 7 Summary and future outlook The focus of this thesis was to explore how GB character, including GB misorientation and GB plane orientation, could affect material behavior such as thermal stability, corrosion resistance, and sensitization. The objective of GBE is to incorporate special GBs into materials for the purpose of improving the properties. Understanding the relationship between GB character and material properties allows for the development of new GBE materials with enhanced properties. Particularly, highly nt Cu, which had bundles of nanoscale-spacing CTBs, was sputtered and tested in order to evaluate the effect of nt GBs (or bundles of CTBs) on the thermal and corrosion behavior. In the thermal study, highly nt Cu was shown to be stable up to 300 C, while non-nt Cu became unstable at 200 C. Furthermore, the polarization and immersion corrosion studies indicated that highly nt Cu had a lower corrosion rate and a more protective passive layer compared with partially nt Cu, non-nt Cu, and mc Cu. Thus, both the thermal and corrosion studies demonstrate the enhancement of properties due to the presence of nanotwins. Further analysis of the nt microstructure revealed that nanotwins could improve the material properties in three major ways. First, nanotwins are thermally stable and resistant to corrosion. Because nanotwins consist of many CTBs which have ordered low-energy GB structures, the nanotwins are much more stable in both elevated- temperature and corrosive environments compared with ordinary GBs. Furthermore, nanotwins form triple junctions with other GBs, which could prohibit grain growth and corrosion propagation. The triple junctions may constrain elevated-temperature GB 121 migration and thus enhance the thermal stability. At each triple junction, GBs may alternate between low-angle and high-angle segments, so the intergranular corrosion could be arrested at the triple junctions by an nt GB and a low-angle GB. In addition to the ordered GB structures and triple junctions, the third effect of nanotwins is that they help to maintain a {111} texture, which may improve thermal stability and corrosion resistance. The orientation of nt GB planes is {111}, and are parallel to the top surface. This leads to a {111} texture on the top surface of highly nt Cu. The {111} texture increases the fraction of low-angle GBs, which then stabilizes the microstructures at elevated temperatures. The {111} crystallographic plane of fcc metals has the highest atomic binding energy, and thus could be more resistant to dissolution during corrosion. A protective passive layer was also observed to be preferentially formed on a {111} textured Cu. Besides nanotwins and their effects on the properties of Cu, Al 5456 (as-received and as-sputtered) samples were studied to explore the role of GB character in the sensitization behavior. The results showed that intergranular β precipitation not only depends on GB misorientation, but that GB plane orientation may also play an important role. Specifically, GB planes with orientations close to {110} were observed to be susceptible to β precipitation. It was also found that AS-5456 had a much better resistance to β precipitation compared with AR-5456, which may be due to its special columnar GB plane orientations. 122 In conclusion, this work has revealed the role of nanotwins in altering the microstructure and subsequently improving the thermal and corrosion properties. It also demonstrates the effect of GB plane orientation on the β precipitation in the Al-Mg alloy. Nevertheless, with more known about the role of GB character in material properties, this study also brings forth more future research directions: (1) While the highly nt Cu in this work had a {111} texture, it is also possible to synthesize highly nt Cu which has a different texture (e.g. {110}, {100}). It will be interesting to evaluate the properties of highly nt Cu with different textures, so that the respective roles of nanotwins and textures could be clarified. (2) Twin thickness is another interesting factor to be explored, because it is directly related to the amount of triple junctions. On one hand, decreasing twin thickness brings more triple junctions, which results in higher resistance to GB migration and frequent alteration of GB characters. On the other hand, decreasing twin thickness increases the overall GB energy, leading to a less stable microstructure. Thus, there might be an optimum twin thickness. (3) Although nanotwins can alternate a high-angle GB to a low-angle GB, some columnar GBs in highly nt Cu were still totally high-angle GBs which are vulnerable to corrosion. So it would be meaningful to synthesize highly nt Cu which has exclusively low-angle or alternating columnar GBs, and assess its thermal and corrosion properties to see if improvement can be achieved by eliminating high-angle columnar GBs. 123 (4) Since AS-5456 showed excellent resistance to β precipitation, it is worthy to further explore the inner mechanism for it. Specifically, one may identify its columnar GB plane orientations, and try to find out if AS-5456 has some preferred GB plane orientations that are immune to β precipitation. It is further suggested that the four experimental approaches mentioned above may be effectively assisted by implementing a modeling approach. These efforts would provide a clearer understanding in the effect of nanotwins and general GB characters on the material properties, so that more GBE materials with enhanced properties can be synthesized and used for industrial applications. 124 Appendix A Abbreviations AR As-received AS As-sputtered bcc Body centered cubic C. D. Current density cg Coarse-grained CSL Coincidence site lattice CTB Coherent twin boundary EBSD Electron backscatter diffraction EDS Energy dispersive spectroscopy EELS Electron energy loss spectroscopy fcc Face centered cubic FIB Focused ion beam GB Grain boundary GBE Grain boundary engineering mc Microcrystalline nc Nanocrystalline ND Normal direction nt Nanotwinned R. T. Room temperature 125 SCE Saturated calomel electrode SEM Scanning electron microscope SFE Stacking fault energy SHE Standard hydrogen electrode STEM Scanning-TEM TEM Transmission electron microscope ufg Ultra-fine grain VD Vertical direction XRD X-ray diffraction 126 Appendix B Further analysis of the anodic Tafel region Fig.71 is the annotated version of Fig.59 in Section 5.3.1, and it shows the anodic and cathodic Tafel linear regions marked as β a and β c , respectively. The interesting thing is that there are transitional regions, as illustrated by the three arrows, from the corrosion potential (E corr ’) to the andic Tafel regions (β a ) for highly nt, partially nt, and mc Cu samples. Figure 71 The annotated version of Fig.59: potentiodynamic polarization curves of highly nt, partially nt, non-nt, and mc Cu in the 3.5% NaCl (pH ~ 8.0) solution. These transitional regions do not seem to be passive regions due to the following reasons. It was stated in Section 5.1 that the potentiodynamic polarization started 127 immediately after the sample was placed into the solution, which means that the corrosion potentials (E corr ’) in Fig.71 were not stabilized and may not represent the actual stable corrosion potentials of the Cu samples. Thus, these unstable corrosion potentials may be changing during potentiodynamic polarization, which could lead to the transitional regions observed in Fig.71. In addition, the relevant literature [141, 152, 153] reported that the initial passive region is usually at potentials above -0.1 V vs. Ag/AgCl with passivation current density (I pass ) typically bigger than 1 mA/cm 2 , which corresponds to the passive region marked by I pass in Fig.71. Therefore, the observed transitional regions in the anodic branch could be due to the unstabilized corrosion potential instead of the passive behavior of Cu samples, and the initial passive region in Fig.71 is expected to be the region indicated by I pass . 128 Appendix C Measurement of grain boundary plane orientations The GB plane orientations presented in Section 6.3 were obtained by analyzing the top-surface EBSD orientation maps and the cross-sectional TEM images. Detailed measurement procedures are explained in this appendix. Figs.72 (a) and (b) illustrate the EBSD top-surface and TEM cross-sectional views of GB-1. The normal direction (ND) marks the crystallographic orientation that is perpendicular to the top surface of the sample, while the vertical direction (VD) represents the crystallographic direction that is in the top surface and points upward in the EBSD map [Fig.72 (a)]. The green arrows in Fig.72 (b) indicate the GB plane orientations of the p1 and p2 segments of GB-1. Figure 72 Top-surface EBSD and cross-sectional TEM views of GB-1 adapted from Fig.68 (b) and Fig.69 (a). ND and VD are normal and vertical directions. The green arrows in (b) represent the GB plane orientations of p1 and p2 segments. 129 VD and ND of grain 1 and grain 2 are known from the EBSD result and listed in Tables 13 – 16. Each row of the data in Tables 13 – 16 represents one EBSD scan point sampled near GB-1 in the respective grain (grain 1 or grain 2). The GB plane orientations of p1 and p2 as illustrated in Fig.72 (b) can then be calculated by counterclockwise rotating ND around VD for 59.7  and 21.5 , respectively. Tables 13 – 16 show the calculated GB plane orientations of p1 and p2 with respect to grain 1 and grain 2. These orientations were unified, averaged, and finally presented in an integer format. Tables 13 and 14 present that p1 has a <10 5 1> orientation with respect to grain 1 and a <7 6 1> orientation with respect to grain 2; Tables 15 and 16 indicate that p2 has a <12 10 9> orientation with respect to grain 1 and a <9 3 2> orientation with respect to grain 2. Table 13 The GB plane orientation of p1 in GB-1 with respect to grain 1. ND and VD are the directions shown in Fig.72. The orientation of p1 is obtained by a 59.7  counterclockwise rotation of ND around VD. Grain 1 ND (h k l) VD [u v w] p1 (59.7  rotation) p1 (unit vector) (-8 10 17) [13 24 -8] [-18.8 9.8 -1.2] <0.885 0.462 0.057> (-13 -6 -8) [4 6 -11] [0.9 -14.5 -7.6] <0.884 0.463 0.055> (11 10 19) [23 -12 -7] [10.6 21.6 -2] <0.895 0.439 0.083> (-18 -13 24) [-11 6 -5] [-14.1 -29.2 -4] <0.894 0.432 0.122> (-12 -6 -7) [2 3 -6] [1 -13.6 -6.5] <0.900 0.430 0.066> (-15 -8 -10) [6 10 -17] [2.3 -17.2 -9.3] <0.874 0.472 0.117> (7 6 12) [6 -3 -2] [6.5 13.6 -1] <0.900 0.430 0.066> (8 16 9) [-8 -5 16] [18 -1.2 8.6] <0.901 0.430 0.060> (11 10 19) [23 -12 -7] [10.6 21.6 -2] <0.895 0.439 0.083> (11 8 18) [24 -15 -8] [11.6 19.3 -1.4] <0.855 0.514 0.062> (-26 -13 -15) [4 7 -13] [2.3 -29 -14.9] <0.887 0.456 0.070> (-23 -11 -13) [4 7 -13] [1.6 -25.4 -13.2] <0.886 0.460 0.056> 130 Table 13 (continued) (9 19 13) [-3 -2 5] [21.5 -2.2 12] <0.870 0.485 0.089> (-17 -23 -15) [-25 10 13] [-12.9 5.6 -29.1] <0.900 0.399 0.173> (7 6 12) [6 -3 -2] [6.5 13.6 -1] <0.900 0.430 0.066> Average <0.888 0.449 0.082> Integer format <10 5 1> Table 14 The GB plane orientation of p1 in GB-1 with respect to grain 2. ND and VD are the directions shown in Fig.72. The orientation of p1 is obtained by a 59.7  counterclockwise rotation of ND around VD. Grain 2 ND (h k l) VD [u v w] p1 (59.7  rotation) p1 (unit vector) (-7 -3 -17) [-13 2 5] [-2.4 14.2 -11.8] <0.763 0.634 0.129> (-6 10 -27) [1 6 2] [21.5 3 -19.8] <0.732 0.674 0.102> (-1 -8 -3) [2 5 -14] [6.8 -5.2 -0.9] <0.790 0.604 0.105> (3 -1 9) [25 3 -8] [0.9 7.6 5.7] <0.796 0.597 0.094> (-3 -1 -7) [-5 1 2] [-0.7 6 -4.8] <0.778 0.622 0.091> (-3 20 8) [4 -9 24] [16.9 13.6 2.3] <0.775 0.623 0.105> (-4 -27 -11) [2 5 -13] [22.9 -18.2 -3.5] <0.777 0.618 0.119> (-3 -1 -7) [-5 1 2] [-0.7 6 -4.8] <0.778 0.622 0.091> (3 -1 9) [25 3 -8] [0.9 7.6 5.7] <0.796 0.597 0.094> (10 -5 26) [14 2 -5] [3.5 21.3 18.3] <0.753 0.647 0.124> (-7 -29 -10) [2 4 -13] [22.7 -21.6 -3.2] <0.721 0.686 0.102> (-2 -1 -5) [-19 3 7] [-0.7 4.1 -3.6] <0.745 0.654 0.127> (11 -6 30) [6 1 -2] [3.1 24.2 21.5] <0.744 0.661 0.095> (11 -5 27) [21 3 -8] [4 22.4 18.9] <0.757 0.639 0.135> (-11 25 -5) [-5 -2 1] [-3.2 18.3 20.6] <0.743 0.660 0.115> Average <0.763 0.636 0.109> Integer format <7 6 1> Table 15 The GB plane orientation of p2 in GB-1 with respect to grain 1. ND and VD are the directions shown in Fig.72. The orientation of p2 is obtained by a 21.5  counterclockwise rotation of ND around VD. Grain 1 ND (h k l) VD [u v w] p2 (21.5  rotation) p2 (unit vector) (-8 10 17) [13 24 -8] [-13.7 11.3 11.7] <0.644 0.550 0.531> 131 Table 15 (continued) (-13 -6 -8) [4 6 -11] [-8.9 -10.5 -8.9] <0.641 0.543 0.543> (11 10 19) [23 -12 -7] [12.4 16.3 12.7] <0.676 0.527 0.515> (-18 -13 24) [-11 6 -5] [-18.9 -21.7 15.5] <0.664 0.578 0.474> (-12 -6 -7) [2 3 -6] [-8.2 -10.1 -7.8] <0.666 0.541 0.514> (-15 -8 -10) [6 10 -17] [-9.8 -13 -11.1] <0.660 0.563 0.497> (7 6 12) [6 -3 -2] [7.8 10.1 8.2] <0.666 0.541 0.514> (8 16 9) [-8 -5 16] [13.4 10.9 10.1] <0.670 0.545 0.505> (11 10 19) [23 -12 -7] [12.4 16.3 12.7] <0.676 0.527 0.515> (11 8 18) [24 -15 -8] [12.8 13.9 12.3] <0.617 0.568 0.546> (-26 -13 -15) [4 7 -13] [-17.6 -21.6 -17.1] <0.661 0.538 0.523> (-23 -11 -13) [4 7 -13] [-15.8 -18.6 -14.9] <0.650 0.553 0.521> (9 19 13) [-3 -2 5] [15.6 12.7 14.4] <0.631 0.582 0.513> (-17 -23 -15) [-25 10 13] [-17.6 -14.1 -23.1] <0.716 0.545 0.437> (7 6 12) [6 -3 -2] [7.8 10.1 8.2] <0.666 0.541 0.514> Average <0.660 0.549 0.511> Integer format <12 10 9> Table 16 The GB plane orientation of p2 in GB-1 with respect to grain 2. ND and VD are the directions shown in Fig.72. The orientation of p2 is obtained by a 21.5  counterclockwise rotation of ND around VD. Grain 2 ND (h k l) VD [u v w] p2 (21.5  rotation) p2 (unit vector) (-7 -3 -17) [-13 2 5] [-6 3.9 -17.2] <0.923 0.322 0.209> (-6 10 -27) [1 6 2] [4.8 8.4 -27.8] <0.944 0.285 0.163> (-1 -8 -3) [2 5 -14] [2.2 -7.9 -2.5] <0.921 0.292 0.257> (3 -1 9) [25 3 -8] [2.5 2.5 8.8] <0.928 0.264 0.264> (-3 -1 -7) [-5 1 2] [-2.5 1.8 -7] <0.915 0.327 0.235> (-3 20 8) [4 -9 24] [5 20.1 6.7] <0.923 0.308 0.230> (-4 -27 -11) [2 5 -13] [6.9 -27 -9.3] <0.919 0.317 0.235> (-3 -1 -7) [-5 1 2] [-2.5 1.8 -7] <0.915 0.327 0.235> (3 -1 9) [25 3 -8] [2.5 2.5 8.8] <0.928 0.264 0.264> (10 -5 26) [14 2 -5] [8.6 5.5 26.4] <0.933 0.304 0.194> (-7 -29 -10) [2 4 -13] [4.6 -29.9 -8.5] <0.952 0.271 0.146> (-2 -1 -5) [-19 3 7] [-1.7 1 -5.1] <0.933 0.311 0.183> (11 -6 30) [6 1 -2] [9.2 6 30.6] <0.941 0.283 0.185> (11 -5 27) [21 3 -8] [9.6 5.9 27.4] <0.925 0.324 0.199> (-11 25 -5) [-5 -2 1] [-9.2 25.7 5.2] <0.925 0.331 0.187> Average <0.928 0.302 0.212> Integer format <9 3 2> 132 The GB plane orientations of p3 and p4 in GB-2 were measured in the same way as p1 and p2 demonstrated above except that the VD needs to be modified to VD’. Figs.73 (a) and (b) show the EBSD top-surface and TEM cross-sectional views of GB-2. Figure 73 Top-surface EBSD and cross-sectional TEM views of GB-2 adapted from Fig.68 (b) and Fig.69 (b). ND and VD are the normal and vertical directions, respectively. The blue arrow in (a) is the VD’ which is parallel to the GB-2 trace and perpendicular to the FIB lift-out as shown by the black rectangle. The green arrows in (b) represent the GB plane orientations of p3 and p4 segments. In Fig.72 (a), VD is parallel to the trace of GB-1, but the VD in Fig.73 (a) is not parallel to the trace of GB-2. Therefore, VD’, which is tilted 74.5  from VD as shown by the blue arrow in Fig.73 (a), is measured and used in the subsequent calculation of p3 and p4 plane orientations. Tables 17 and 18 list the detailed calculation procedures of p3 133 plane orientations with respect to grain 3 and grain 4. Fig.73 (b) shows that the GB plane orientation of p3 is 11.7  clockwise rotation of ND around VD’. After unifying and averaging the p3 orientations, Tables 17 and 18 indicate that the p3 GB plane has a <21 16 1> orientation with respect to grain 3 and a <9 8 2> orientation with respect to grain 4. Follow similar measurement procedures, Tables 19 and 20 show that the p4 GB plane has a <14 7 1> orientation with respect to grain 3 and a <7 4 1> orientation with respect to grain 4. Table 17 The GB plane orientation of p3 in GB-2 with respect to grain 3. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB-2 and tilted 74.5  from VD. The orientation of p3 is obtained by a 11.7  clockwise rotation of ND around VD’. Grain 3 ND (h k l) VD [u v w] VD’ [x y z] (74.5 ) p3 (11.7 , unit vector) (-18 -16 -1) [-17 19 2] [-4 3 25.1] <0.800 0.601 0> (-20 1 23) [8 -1 7] [1.2 -10.5 1.5] <0.793 0.609 0.003> (-8 7 0) [-7 -8 -1] [-1.2 -1.4 -10.5] <0.798 0.601 0.038> (-25 21 1) [-11 -13 -2] [-2.1 -1.7 -16.9] <0.784 0.621 0> (-6 5 0) [-15 -18 -2] [-2.8 -3.3 -23.1] <0.779 0.626 0.038> (-8 -7 0) [-7 8 1] [-1.2 1.4 10.5] <0.798 0.601 0.038> (0 15 17) [-2 -17 15] [-22.4 -3.1 2.7] <0.797 0.603 0.035> (-1 17 -20) [1 -7 -6] [9.1 -0.9 -1.2] <0.787 0.616 0.004> (-8 7 0) [-7 -8 -1] [-1.2 -1.4 -10.5] <0.798 0.601 0.038> (0 16 -19) [3 -19 -16] [24.7 -2.9 -2.4] <0.784 0.619 0.032> (-1 17 -19) [1 -10 -9] [13.2 -1.6 -2.1] <0.803 0.596 0> Average <0.793 0.609 0.021> Integer format <21 16 1> 134 Table 18 The GB plane orientation of p3 in GB-2 with respect to grain 4. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB-2 and tilted 74.5  from VD. The orientation of p3 is obtained by a 11.7  clockwise rotation of ND around VD’. Grain 4 ND (h k l) VD [u v w] VD’ [x y z] (74.5 ) p3 (11.7 , unit vector) (-6 16 -23) [3 -19 -14] [23 0.1 -6] <0.715 0.676 0.178> (-14 3 17) [19 -2 16] [1.5 -24.3 5.5] <0.774 0.626 0.095> (-4 1 5) [23 -3 19] [1.1 -29.2 6.7] <0.761 0.637 0.124> (5 -15 -21) [-3 20 -15] [-24.4 0.3 -6] <0.726 0.669 0.156> (5 15 22) [-1 -7 5] [-8.4 -0.2 2] <0.712 0.686 0.151> (1 3 4) [-3 -23 18] [-28.4 -0.5 7.5] <0.746 0.648 0.157> (-19 6 26) [22 -4 17] [-0.2 -27.4 6.2] <0.733 0.662 0.156> (-1 3 -4) [2 -14 -11] [17.4 -0.2 -4.5] <0.746 0.648 0.157> (5 -15 -21) [-3 20 -15] [-24.4 0.3 -6] <0.726 0.669 0.156> (-4 1 5) [23 -3 19] [1.1 -29.2 6.7] <0.761 0.637 0.124> (-1 3 -4) [3 -23 -18] [28.4 -0.5 -7.5] <0.746 0.648 0.157> (-6 16 -23) [3 -19 -14] [23 0.1 -6] <0.715 0.676 0.178> (-2 5 -7) [1 -8 -6] [9.7 -0.1 -2.8] <0.722 0.665 0.192> (-5 14 -19) [3 -22 -17] [27 -0.2 -7.3] <0.733 0.659 0.170> (5 16 21) [-4 -25 20] [-31.4 -0.1 7.5] <0.746 0.649 0.153> Average <0.737 0.657 0.154> Integer format <9 8 2> Table 19 The GB plane orientation of p4 in GB-2 with respect to grain 3. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB-2 and tilted 74.5  from VD. The orientation of p4 is obtained by a 23.4  clockwise rotation of ND around VD’. Grain 3 ND (h k l) VD [u v w] VD’ [x y z] (74.5 ) p4 (23.4 , unit vector) (-18 -16 -1) [-17 19 2] [-4 3 25.1] <0.904 0.427 0.037> (-20 1 23) [8 -1 7] [1.2 -10.5 1.5] <0.899 0.436 0.039> (-8 7 0) [-7 -8 -1] [-1.2 -1.4 -10.5] <0.900 0.431 0.066> (-25 21 1) [-11 -13 -2] [-2.1 -1.7 -16.9] <0.892 0.451 0.034> (-6 5 0) [-15 -18 -2] [-2.8 -3.3 -23.1] <0.884 0.461 0.077> (-8 -7 0) [-7 8 1] [-1.2 1.4 10.5] <0.900 0.431 0.066> 135 Table 19 (continued) (0 15 17) [-2 -17 15] [-22.4 -3.1 2.7] <0.901 0.428 0.071> (-1 17 -20) [1 -7 -6] [9.1 -0.9 -1.2] <0.895 0.446 0.027> (-8 7 0) [-7 -8 -1] [-1.2 -1.4 -10.5] <0.900 0.431 0.066> (0 16 -19) [3 -19 -16] [24.7 -2.9 -2.4] <0.890 0.451 0.060> (-1 17 -19) [1 -10 -9] [13.2 -1.6 -2.1] <0.904 0.426 0.039> Average <0.897 0.438 0.053> Integer format <14 7 1> Table 20 The GB plane orientation of p4 in GB-2 with respect to grain 4. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB-2 and tilted 74.5  from VD. The orientation of p4 is obtained by a 23.4  clockwise rotation of ND around VD’. Grain 4 ND (h k l) VD [u v w] VD’ [x y z] (285.5 ) p4 (23.4 , unit vector) (-6 16 -23) [3 -19 -14] [23 0.1 -6] <0.841 0.524 0.136> (-14 3 17) [19 -2 16] [1.5 -24.3 5.5] <0.885 0.463 0.049> (-4 1 5) [23 -3 19] [1.1 -29.2 6.7] <0.876 0.476 0.077> (5 -15 -21) [-3 20 -15] [-24.4 0.3 -6] <0.849 0.514 0.118> (5 15 22) [-1 -7 5] [-8.4 -0.2 2] <0.839 0.532 0.111> (1 3 4) [-3 -23 18] [-28.4 -0.5 7.5] <0.863 0.491 0.118> (-19 6 26) [22 -4 17] [-0.2 -27.4 6.2] <0.855 0.504 0.119> (-1 3 -4) [2 -14 -11] [17.4 -0.2 -4.5] <0.863 0.491 0.118> (5 -15 -21) [-3 20 -15] [-24.4 0.3 -6] <0.849 0.514 0.118> (-4 1 5) [23 -3 19] [1.1 -29.2 6.7] <0.876 0.476 0.077> (-1 3 -4) [3 -23 -18] [28.4 -0.5 -7.5] <0.863 0.491 0.118> (-6 16 -23) [3 -19 -14] [23 0.1 -6] <0.841 0.524 0.136> (-2 5 -7) [1 -8 -6] [9.7 -0.1 -2.8] <0.848 0.509 0.147> (-5 14 -19) [3 -22 -17] [27 -0.2 -7.3] <0.856 0.500 0.128> (5 16 21) [-4 -25 20] [-31.4 -0.1 7.5] <0.866 0.487 0.115> Average <0.858 0.500 0.112> Integer format <7 4 1> 136 Appendix D List of figures and tables Figures Figure 1 Atom positions of three crystal structures: bcc, fcc, hcp [25]. ...........................5 Figure 2 3D reconstruction of Ni-based superalloy by EBSD [26]..................................6 Figure 3 The soap bubble model of low-angle GB and high-angle GB [27]. ...................7 Figure 4 Schematic graph of a 5 CSL: red and yellow dots represent the atoms of two crystal lattices, and green dots represent the coincident sites [28]. ....................................9 Figure 5 The computed GB energies of Ni for ranging from 3 to 400.  is the inverse density of coincident sites. Different symbols in the inset represent different types of GBs [7]. ......................................................................................................................... 10 Figure 6 Schematic atomic arrangement on the two sides of a CTB in fcc metals: (a) shows the {111} planes marked as “A”, “B”, “C” (vertical direction), adapted from [29]; (b) presents the hcp structure of the CTB, adapted from [30]. ........................................ 11 Figure 7 GB distributions of the (a) starting base stainless steel, and (b) GBE stainless steel. Black and gray lines represent ordinary and special GBs, respectively [33]. ......... 12 Figure 8 Cross-sectional optical micrographs of (a) conventional and (b) GBE alloy 800 (Fe-35Ni-25Cr, heat-treated at 600°C for 1 hour) after 120-hours exposure to solution containing sulfuric acid and ferric sulfate (ASTM G28 standard) [3]. ............................ 13 Figure 9 Cross-sectional models and optical micrographs showing the propagation of intergranular corrosion through the GB network. In (a) and (b), “R” refers to random or ordinary GBs; “c” and “i” refer to coherent and incoherent 3 GBs, respectively. The thick black lines are corroded GBs, and the thin lines are non-corroded GBs as shown by the legend at bottom right. (a) and (b) are from ref. [19]; (c) is from ref. [34]. ............... 14 Figure 10 Atomic arrangement of nanotwins studied by high-resolution TEM. The solid white lines indicate (1 1 ¯ 1), (1 ¯ 1 ¯ 1), and (2 0 0) planes, respectively [35]. ...................... 16 Figure 11 Schematic of the magnetron sputtering process, adapted from [41]. ............. 17 137 Figure 12 TEM micrographs of the (a) top surface, (b) cross-section of magnetron sputtered nt Cu [42], and the (c) top surface of electrodeposited nt Cu [43]. ................... 18 Figure 13 TEM grain orientation mappings of columnar GBs intersected by nanotwins: (a) an alternating columnar GB, and (b) a high-angle columnar GB, indicated by the two black arrows. GB types are as follows: black, high-angle GB; yellow, low-angle GB; red, 3 GB; blue, 9 GB. In (b), the white arrow refers to the growth direction of the foil, and the inset correlates each color with a crystal orientation [50]. .................................. 19 Figure 14 Tensile true stress – true strain curves of as-deposited nt Cu (average grain size, d avg  400 nm) in comparison with nc Cu (d avg  30 nm) and cg Cu (d avg > 100 µm) [43]................................................................................................................................ 21 Figure 15 Tensile true stress-true strain curves of nt Cu samples (400 nm < d avg < 600 nm) with various twin thicknesses (from 96 nm down to 4 nm) in comparison with twin- free ufg Cu (d avg  500 nm) and cg Cu (d avg  10 µm) [48]. ........................................... 22 Figure 16 Grain growth at different heat-treatment temperatures for nt Cu (red circles) and ufg Cu (blue triangles). Red and blue lines are the fitted grain growth curves calculated from normal grain growth equation. Adapted from [8]. ................................ 25 Figure 17 Cross-sectional TEM micrographs of as-sputtered nt Cu (a) before and (b) after heat-treatment at 800°C for one hour [8, 9]. ........................................................... 25 Figure 18 Potentiodynamic polarization curves of nc (d avg  32 nm) and cg (d avg  100 m) 99.99% nickel using a scan rate of 0.5 mV/s in 1 mol/L H 2 SO 4 at 293K [67]. ......... 29 Figure 19 Potentiodynamic polarization curves of nc (grain size, 56 nm) and mc (grain size, 2 m) Cu coatings using a scan rate of 0.5 mV/s in 0.1 mol/L NaOH solution [73]. ...................................................................................................................................... 30 Figure 20 Potentiodynamic polarization curves of nt Ni and cast Ni in 0.1 mol/L H 3 BO 3 + 0.025 mol/L Na 2 B 4 O 7 solution [15]. ........................................................................... 31 Figure 21 Al-Mg phase diagram, adapted from [95]. .................................................... 33 Figure 22 Dark field TEM images of (a) continuous and (b) discontinuous intergranular β phase, adapted from [2]............................................................................................... 34 138 Figure 23 Normalized strain to failure versus the time of heat-treatment (at 150 C) plot obtained by the constant extension rate testings of Al-5083 samples in a 3.5% NaCl solution. Insets are the TEM images showing β precipitation at selected heat-treatment times. Adapted from [2]. ............................................................................................... 35 Figure 24 The (a) model and (b) simplified model (originated by Zener [98]) of solute concentration profile during diffusion controlled β phase growth [97]. C β , C 0 , and C e refer to the Mg concentrations in β phase, Al-Mg matrix, and equilibrium state. ............ 37 Figure 25 The relationship between GB misorientation angle and attacked mode obtained from etching a Al 5182 sample sensitized at 150 C for 10 hours [22]. ............. 38 Figure 26 (a) Bright-field image of an intergranular β phase in a sensitized Al 5182; (b) magnified image of a β phase with the normal directions (N 1 , N 2 ) of its interfacial planes. Adapted from [101]. ...................................................................................................... 40 Figure 27 Schematic of Bragg’s diffraction by crystal planes [103]. ............................. 42 Figure 28 JSM-7001F-LV field-emission SEM equipped with an EBSD analytical detector. ......................................................................................................................... 44 Figure 29 The electron – specimen atom interaction volume under the specimen surface [102]. ............................................................................................................................. 45 Figure 30 Schematic set-up of a conventional EBSD system [107]. ............................. 46 Figure 31 (a) Schematic of the interaction between electron beam and tilted specimen, and the subsequent Kikuchi line generation; (b) A Kikuchi pattern from cadmium at 20 keV [108]. ..................................................................................................................... 47 Figure 32 Schematic of EBSD grain orientation image. Grain orientations are presented by the colored crystal lattices [109]. ............................................................................... 47 Figure 33 GB character distribution of a GBE brass. Black, red, and blue lines represent ordinary, 3, and 9 GBs respectively. Adapted from [5]. ............................................ 48 139 Figure 34 The experimental set-up of transmission EBSD. The dark blue, red, and light blue lines refer to the incident electron beam, plane of the thin film, and the transmitted electrons towards the EBSD detector [112]. ................................................................... 49 Figure 35 Schematic illustration of the interaction between an incident Ga + ion and crystal lattice atoms. R p and R l are the projected and lateral ranges of the implanted Ga + ion [115]. ....................................................................................................................... 51 Figure 36 Schematic images illustrating the effect of (a), (b) grain orientation, and (c) atomic mass on the emitted secondary electrons. The orange atoms in (c) have bigger atomic mass compared with those in (a), (b). Adapted from [115]. ................................ 53 Figure 37 (a) The lift-out of an area of interest from the sample by a probe tip; (b) The attachment of the area of interest to a TEM grid for thinning and TEM observation. ...... 54 Figure 38 TEM image formation by diffraction contrast: (a) detailed formation process; (b) bright-field image formation; (c) dark-field image formation. Adapted from [102]. . 55 Figure 39 (a) Inelastic scattering of the incident electron on the specimen atom; (b) EELS energy loss spectrum of boron nitride (BN). Adapted from [117]. ....................... 56 Figure 40 MTI-GSL1100X vacuum furnace system. .................................................... 57 Figure 41 Temperature profiles of the heat-treating, holding and cooling processes used in this study. .................................................................................................................. 58 Figure 42 The (a) actual and (b) schematic electrochemical cell showing the three electrodes (black clip – working electrode; yellow clip – counter electrode). (c) Epoxy- mounted sample as the working electrode, and (d) the Gamry Reference 3000 potentiostat. ................................................................................................................... 59 Figure 43 Schematic plot of a linear polarization curve [120]. ...................................... 61 Figure 44 Schematic of the potentiodynamic polarization curve for a passive material. “C. D.” stands for current density. Adapted from [120]. ................................................ 62 140 Figure 45 The (a) optical, (b) SEM, and (c) EBSD surface micrographs of the as- received Al 5456 after 10% H 3 PO 4 etching at 50 C for 2 minutes. The triangle legend correlates the colors in (c) with specific grain orientations. ............................................ 66 Figure 46 The optical surface graphs of as-received Al 5456 and sensitized Al 5456 before (a, b) and after (c, d) etching in 10% H 3 PO 4 at 35 C for 1 min. ........................... 67 Figure 47 EBSD scan area and its near-by fiducial nanoindents in the thermal study. ... 69 Figure 48 FIB micrographs of as-prepared ufg Cu on (a) top surface, (b) cross-section, and nt Cu on (c) top surface, (d) cross-section. The arrows in (b) and (d) indicate the growth direction of foils. ............................................................................................... 71 Figure 49 EBSD top-surface grain morphologies of ufg Cu (a) as-prepared, and after (b) 200°C, (c) 300°C, (d) 400°C heat-treatments. Each color represents a grain, not a grain orientation. .................................................................................................................... 72 Figure 50 EBSD top-surface grain morphologies of nt Cu (a) as-prepared, and after (b) 200°C, (c) 300°C, (d) 400°C heat-treatments. Each color represents a grain, not a grain orientation. .................................................................................................................... 73 Figure 51 EBSD top-surface grain size distribution of nt Cu (a1) as-prepared (average grain size d = 697 nm), and after (a2) 200°C (d = 734 nm), (a3) 300°C (d = 754 nm), (a4) 400°C (d = 1195 nm) heat-treatments; and those of ufg Cu (b1) as-prepared (d = 328 nm), and after (b2) 200°C (d = 963 nm), (b3) 300°C (d = 1871 nm), (b4) 400°C (d = 4364 nm) heat-treatments. The insets represent size distributions of bigger grains. ....................... 74 Figure 52 FIB micrographs after 400°C heat-treatment of ufg Cu on (a) top surface, (b) cross-section, and nt Cu on (c) top surface, (d) cross-section. The arrows in (b) and (d) indicate the growth direction of foils. ............................................................................. 75 Figure 53 Top-surface GB distributions (EBSD) of ufg Cu (a) as-prepared, (b) after 400°C heat-treatment, and nt Cu (c) as-prepared, (d) after 400°C heat-treatment. Red lines represent 3 GBs; green lines represent low-angle GBs (θ = 5 – 15°), and black lines represent high angle GBs (but not 3, θ = 15 – 63°). ............................................. 76 Figure 54 XRD peaks of (a) ufg Cu and (b) nt Cu before and after each heat-treatment. “R.T.” refers to room temperature. ................................................................................. 79 141 Figure 55 Top-surface EBSD inverse pole figure maps of ufg Cu (a1) as-prepared, and after (a2) 200°C, (a3) 300°C, (a4) 400°C heat-treatments; nt Cu (b1) as-prepared, and after (b2) 200°C, (b3) 300°C, (a4) 400°C heat-treatments. The inset of each map is the inverse pole triangle, and the triangle at the bottom right indicates the crystallographic orientation of the maps and insets. ................................................................................. 81 Figure 56 Top-surface (a) FIB micrograph and (b) XRD peaks of nt Cu after heat- treatment at 600°C. ........................................................................................................ 83 Figure 57 FIB cross-sectional micrographs of as-prepared (a) highly nt, (b) partially nt, (c) non-nt, and (d) mc Cu samples. The arrows indicate growth direction and point to the top surfaces which will be exposed to the corrosion solution. ......................................... 89 Figure 58 EBSD top-surface grain orientation map overlapped with GBs: as-prepared (a) highly nt, (b) partially nt, (c) non-nt, (d) mc Cu. The inset in each map is the inverse pole triangle, and the triangle legend at bottom right indicates specific crystallographic orientations in the maps and insets. The legend at top right correlates twin boundaries, high-angle and low-angle GBs with red, black, and green lines, respectively. In a, b and c, the observed twin boundaries are mostly due to the nanotwins intersecting with the surface; the curled shape of these twin boundaries is due to the surface roughness. ........ 91 Figure 59 Potentiodynamic polarization curves of highly nt, partially nt, non-nt, and mc Cu in the 3.5% NaCl (pH ~ 8.0) solution. The two y-axes are potentials versus Ag/AgCl and SHE reference potentials. ........................................................................................ 94 Figure 60 (a) corrosion current density (I corr ) and (b) passivation current density (I pass ) of highly nt, partially nt, non-nt, and mc Cu samples in the 3.5% NaCl (pH ~ 8.0) solution. ...................................................................................................................................... 95 Figure 61 Schematic illustration of two possible columnar grain boundary configurations within a nt Cu sample: (a) alternating high-angle and low-angle boundary segments, and (b) high-angle boundary only. The legend at bottom correlates red, black, and green lines with nanotwin, high-angle, and low-angle GBs, respectively. Different crystal orientations are illustrated by different gray-scale colors. The two arrows point along the corrosion surface. ........................................................................................... 97 142 Figure 62 SEM top-surface morphologies and FIB cross-sectional micrographs of passivated (a1, a2) highly nt, (b1, b2) partially nt, (c1, c2) non-nt, and (d1, d2) mc Cu during polarization in the 3.5% NaCl (pH ~ 8.0) solution. Markers “o” and “x” indicate nt and non-nt columnar grains, respectively; each arrow marker points to a specific GB. The black layer on the top of each FIB micrograph is the FIB deposited carbon layer. The actual length in the vertical direction of FIB micrographs is  1.27 times the imaged length due to the 52  tilt. ................................................................................................ 98 Figure 63 SEM top-surface topographies of the passive layers on (a) highly nt, (b) partially nt, (c) non-nt, and (d) mc Cu samples after immersion corrosion in 3.5% NaCl (pH ~ 8.0) solution....................................................................................................... 102 Figure 64 SEM top-surface topographies of the passive layers on (a) highly nt, (b) partially nt, (c) non-nt, and (d) mc Cu samples after immersion corrosion in 3.5% NaCl (pH ~ 8.0) solution and subsequent ultrasonic cleaning. ............................................... 103 Figure 65 Normalized XRD peaks showing the corrosion products of (a) passivated Cu, and (b) immersion corroded and subsequently ultrasonic cleaned Cu samples. The solid squares and circles on top of the peaks mark CuCl and Cu 2 O, respectively; the peak orientations are also marked......................................................................................... 105 Figure 66 Method for the observation of intergranular β phase: top-surface (a) phosphoric etching of the β phase, (b) FIB milling and thinning, and cross-sectional (c) TEM observation of the FIB lift-out foil. Adapted from [124]. .................................... 110 Figure 67 (a) Top-surface EBSD map of AR-5456, and (b) cross-sectional dark-field TEM graphs of AS-5456. The triangle legend in (a) correlates the color of the EBSD map with grain orientations.......................................................................................... 111 Figure 68 Top-surface (a) SEM topography and (b) EBSD grain orientation map of the identical area of the heat-treated/etched AR-5456. The white rectangles in (a) are the FIB lift-out areas, and the circled numbers mark four distinct grains. Red, yellow and black lines in (b) represent low-CSL, low-angle and high-angle GBs, respectively. The bottom- right triangle correlates each color in the EBSD map with a specific grain orientation. 113 143 Figure 69 Cross-sectional TEM micrographs of the two GBs lifted out by FIB in Fig.68: (a) GB-1, and (b) GB-2. Mg composition maps of the dashed box areas were obtained by EELS. The circled numbers mark the four grains, and p1 – p4 indicate the estimated GB plane traces. (c) GB plane orientations of p1 – p4 in an inverse pole triangle: darker dots represent thicker β phase, and hollow dots refer to no β precipitation. .......................... 115 Figure 70 Top-surface (a) SEM and cross-sectional (b) TEM, (c) magnified TEM, and (d) transmission-EBSD graphs of heat-treated/etched AS-5456. In (c), the yellow arrows of GB-IV refer to one Mg-rich region. The grain orientations (normal to the graphic plane) and GB types in (d) are defined by the legends in Fig.68. .................................. 117 Figure 71 The annotated version of Fig.59: potentiodynamic polarization curves of highly nt, partially nt, non-nt, and mc Cu in the 3.5% NaCl (pH ~ 8.0) solution. .......... 126 Figure 72 Top-surface EBSD and cross-sectional TEM views of GB-1 adapted from Fig.68 (b) and Fig.69 (a). ND and VD are normal and vertical directions. The green arrows in (b) represent the GB plane orientations of p1 and p2 segments. .................... 128 Figure 73 Top-surface EBSD and cross-sectional TEM views of GB-2 adapted from Fig.68 (b) and Fig.69 (b). ND and VD are the normal and vertical directions, respectively. The blue arrow in (a) is the VD’ which is parallel to the GB-2 trace and perpendicular to the FIB lift-out as shown by the black rectangle. The green arrows in (b) represent the GB plane orientations of p3 and p4 segments. ................................... 132 Tables Table 1 Misorientation angles and axes of special GBs in cubic-crystal metals [28]. ......8 Table 2 Corrosion parameters derived from the polarization curves of nt Ni and cast Ni in Fig.19 [15]. ................................................................................................................ 31 Table 3 Typical grain orientations{hkl} and the corresponding d and 2θ of Cu (PDF#98- 000-0172), CuCl (PDF#98-000-0323), Cu 2 O (PDF#98-000-0186), and Al (PDF#00-004- 0787). ............................................................................................................................ 43 Table 4 Conversion of potentials among Ag/AgCl, SCE, and SHE. ............................. 60 144 Table 5 β phase etching conditions............................................................................... 65 Table 6 Length fractions (EBSD) of general 3 and coherent 3 GBs on the top surfaces of ufg Cu and nt Cu before and after each heat-treatment. .............................................. 77 Table 7 Area fractions (EBSD) of major orientations on the top surfaces of ufg Cu and nt Cu before and after each heat-treatment (orientations deviated less than 7° from the exact major orientation are counted as this major orientation). ....................................... 82 Table 8 EBSD quantitative information of grain size, grain boundaries, and grain orientations on the top surface of the Cu samples........................................................... 92 Table 9 Summary of polarization tests results. ............................................................. 93 Table 10 Grain sizes and elemental compositions of AR-5456 and AS-5456 before and after heat-treatment. ..................................................................................................... 112 Table 11 Different GB misorientations and the corresponding etching behaviors. ...... 114 Table 12 β phase thicknesses on the different GB planes of the sensitized AR-5456. . 116 Table 13 The GB plane orientation of p1 in GB-1 with respect to grain 1. ND and VD are the directions shown in Fig.72. The orientation of p1 is obtained by a 59.7  counterclockwise rotation of ND around VD................................................................ 129 Table 14 The GB plane orientation of p1 in GB-1 with respect to grain 2. ND and VD are the directions shown in Fig.72. The orientation of p1 is obtained by a 59.7  counterclockwise rotation of ND around VD................................................................ 130 Table 15 The GB plane orientation of p2 in GB-1 with respect to grain 1. ND and VD are the directions shown in Fig.72. The orientation of p2 is obtained by a 21.5  counterclockwise rotation of ND around VD................................................................ 130 Table 16 The GB plane orientation of p2 in GB-1 with respect to grain 2. ND and VD are the directions shown in Fig.72. The orientation of p2 is obtained by a 21.5  counterclockwise rotation of ND around VD................................................................ 131 145 Table 17 The GB plane orientation of p3 in GB-2 with respect to grain 3. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB- 2 and tilted 74.5  from VD. The orientation of p3 is obtained by a 11.7  clockwise rotation of ND around VD’. ......................................................................................... 133 Table 18 The GB plane orientation of p3 in GB-2 with respect to grain 4. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB- 2 and tilted 74.5  from VD. The orientation of p3 is obtained by a 11.7  clockwise rotation of ND around VD’. ......................................................................................... 134 Table 19 The GB plane orientation of p4 in GB-2 with respect to grain 3. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB- 2 and tilted 74.5  from VD. The orientation of p4 is obtained by a 23.4  clockwise rotation of ND around VD’. ......................................................................................... 134 Table 20 The GB plane orientation of p4 in GB-2 with respect to grain 4. ND and VD are the directions shown in Fig.73. VD’ is the direction that is parallel to the trace of GB- 2 and tilted 74.5  from VD. The orientation of p4 is obtained by a 23.4  clockwise rotation of ND around VD’. ......................................................................................... 135 146 References [1] Jones RH, Baer DR, Danielson MJ, Vetrano JS. Metall and Mat Trans A 2001;32A:1699. [2] Searles JL, Gouma PI, Buchheit RG. Metall and Mat Trans A 2001;32A:2859. [3] Palumbo G, Lehockey EM, Lin P. JOM 1998;50:40. [4] Randle V. Scripta Materialia 2006;54:1011. [5] Randle V, Hu Y. The role of vicinal 3 boundaries and 9 boundaries in grain boundary engineering. Eleventh international conference on intergranular and interphase boundaries 2004, vol. 40: Journal of Materials Science, 2005. p.3243 [6] Randle V. Acta Materialia 2004;52:4067. [7] Olmsted DL, Foiles SM, Holm EA. Acta Materialia 2009;57:3694. [8] Anderoglu O, Misra A, Wang H, Zhang X. J. Appl. Phys. 2008;103:094322. [9] Zhang X, Misra A. Scripta Materialia 2012;66:860. [10] Xu D, Sriram V, Ozolins V, Yang JM, Tu KN, Stafford GR, Beauchamp C, Zienert I, Geisler H, Hofmann P, Zschech E. Microelectronic Engineering 2008;85:2155. [11] Saldana C, Murthy TG, Shankar MR, Stach EA, Chandrasekar S. Applied Physics Letters 2009;94:021910. [12] Saldana C, King AH, Stach EA, Compton WD, Chandrasekar S. Applied Physics Letters 2011;99:231911 [13] Murr LE. Interfacial phenomena in metals and alloys. Reading, Massachusetts: Addison-Wesley, 1975. [14] Bernstein N, Tadmor EB. Physical Review B 2004;69:094116. [15] Meng G, Shao Y, Zhang T, Zhang Y, Wang F. Electrochimica Acta 2008;53:5923. [16] Sun F, Meng G, Zhang T, Shao Y, Wang F, Dong C, Li X. Electrochimica Acta 2009;54:1578. 147 [17] Meng G, Sun F, Shao Y, Zhang T, Wang F, Dong C, Li X. Electrochimica Acta 2010;55:2575. [18] Meng G, Wei L, Shao Y, Zhang T, Wang F, Dong C, Li X. JES 2009;156:C240. [19] Hu C, Xia S, Li H, Liu T, Zhou B, Chen W, Wang N. Corrosion Science 2011;53:1880. [20] Ralston KD, Birbilis N. Corrosion 2010;66:075005:1. [21] Mayanna SM, Setty THV. Indian Journal of Chemistry 1972;10:295. [22] Davenport AJ, Yuan Y, Ambat R, Connolly BJ, Strangwood M, Afseth A, Scamans G. Materials Science Forum 2006;519-521:641. [23] Kaigorodova LI. Materials Science Forum 1999;294-296:477. [24] Scotto D’Antuono D, Gaies J, Golumbfskie W, Taheri ML. Scripta Materialia 2014;76:81. [25] Atomic arrangements of three crystal structures. vol. 31.3 KB. http://engineeronadisk.com/notes_material/materiala3.html: Engineerondisk.com, 2014. [26] Groeber M, Ghosh S, Uchic MD, Dimiduk DM. Acta Materialia 2008;56:1257. [27] Gottstein G, Shvindlerman LS. Structure of grain boundaries. In: Ralph B, editor. Grain boundary migration in metals: Thermodynamics, kinetics, applications, second edition. CRC Press, 2010. p.111. [28] EDAX. OIM analysis 6 help manual. AMETEK Inc., 2011. [29] Lu L, Lu K, Suresh S. Science 2009;324:349. [30] Wood EL, Sansoz F. Nanoscale 2012;4:5268. [31] Schuh CA, Kumar M, King WE. Acta Materialia 2003;51:687. [32] Schwartz AJ, King WE, Kumar M. Scripta Materialia 2006;54:963. [33] Michiuchi M, Kokawa H, Wang ZJ, Sato YS, Sakai K. Acta Materialia 2006;54:5179. [34] Shimada M, Kokawa H, Wang ZJ, Sato YS, Karibe I. Acta Materialia 2002;50:2331. 148 [35] Jang D, Li X, Gao H, Greer JR. Nat Nano 2012;advance online publication. [36] Schramm RE, Reed RP. MTA 1975;6:1345. [37] Cai J, Shekhar S, Wang J, Ravi Shankar M. Scripta Materialia 2009;60:599. [38] Yu Q, Qi L, Chen K, Mishra RK, Li J, Minor AM. Nano Letters 2012;12:887. [39] Zhang X, Anderoglu O, Hoagland R, Misra A. JOM 2008;60:75. [40] Hsiao H-Y, Liu C-M, Lin H-w, Liu T-C, Lu C-L, Huang Y-S, Chen C, Tu KN. Science 2012;336:1007. [41] Magnetron sputter coating technology. vol. 2014. http://www.guardian- russia.ru/en/about-glass/modern-technologies/magnetron-sputtering-technology/: Guardian-russia. [42] Shute CJ, Myers BD, Xie S, Li SY, Barbee Jr TW, Hodge AM, Weertman JR. Acta Materialia 2011;59:4569. [43] Lu L, Shen Y, Chen X, Qian L, Lu K. Science 2004;304:422. [44] Hodge AM, Wang YM, Barbee Jr TW. Materials Science and Engineering: A 2006;429:272. [45] Zhang X, Wang H, Chen XH, Lu L, Lu K, Hoagland RG, Misra A. Applied Physics Letters 2006;88:173116. [46] Anderoglu O, Misra A, Wang H, Ronning F, Hundley MF, Zhang X. Applied Physics Letters 2008;93:083108. [47] Ma E, Wang YM, Lu QH, Sui ML, Lu L, Lu K. Applied Physics Letters 2004;85:4932. [48] Lu L, Chen X, Huang X, Lu K. Science 2009;323:607. [49] You ZS, Lu L, Lu K. Acta Materialia 2011;59:6927. [50] Wang YM, Sansoz F, LaGrange T, Ott RT, Marian J, Barbee Jr TW, Hamza AV. Nat Mater 2013;12:697. [51] Shute CJ, Myers BD, Xie S, Barbee Jr TW, Hodge AM, Weertman JR. Scripta Materialia 2009;60:1073. 149 [52] Gottstein G, Shvindlerman LS. Grain boundary motion. In: Ralph B, editor. Grain boundary migration in metals: Thermodynamics, kinetics, applications, second edition. CRC Press, 2010. p.135. [53] Thompson CV, Carel R. Materials Science Forum 1996;204 - 206:83. [54] Sonnweber-Ribic P, Gruber P, Dehm G, Arzt E. Acta Materialia 2006;54:3863. [55] Zhang JM, Xu KW, Ji V. Applied Surface Science 2002;187:60. [56] Ma F, Zhang JM, Xu KW. Applied Surface Science 2005;242:55. [57] Mirpuri K, Wendrock H, Menzel S, Wetzig K, Szpunar J. Thin Solid Films 2006;496:703. [58] Zhang JM, Xu KW, Ji V. Applied Surface Science 2002;185:177. [59] Roberge PR. Corrosion basics. In: Hager LS, editor. Corrosion engineering: principles and practice. McGraw-Hill Professional, 2008. p.19. [60] Roberge PR. The study of corrosion. In: Hager LS, editor. Corrosion engineering: principles and practice. McGraw-Hill Professional, 2008. p.754. [61] Schmitt G. Global needs for knowledge dissemination, research, and development in materials deterioration and corrosion control. World corrosion organization, 2009. p.44. [62] Palumbo G, Aust KT. Acta Metallurgica et Materialia 1990;38:2343. [63] Erb U, Gleiter H, Schwitzgebel G. Acta Metallurgica 1982;30:1377. [64] Lin P, Palumbo G, Erb U, Aust KT. Scripta Metallurgica et Materialia 1995;33:1387. [65] Fang X, Wang W, Guo H, Qin C, Zhou B. International Journal of Modern Physics B 2009;23:1110. [66] Miyamoto H, Ikeuchi K, Mimaki T. Scripta Materialia 2004;50:1417. [67] Rofagha R, Langer R, El-Sherik AM, Erb U, Palumbo G, Aust KT. Scripta Metallurgica et Materialia 1991;25:2867. [68] Pinto EM, Ramos AS, Vieira MT, Brett CMA. Corrosion Science 2010;52:3891. 150 [69] Shi P, Wang Q, Xu Y, Luo W. Materials Letters 2011;65:857. [70] Luo W, Shi P, Wang C, Xu Y, Yan M. Journal of The Electrochemical Society 2012;159:C80. [71] Luo W, Qian C, Wu XJ, Yan M. Materials Science and Engineering: A 2007;452–453:524. [72] Barbucci A, Farne G, Matteazzi P, Riccieri R, Cerisola G. Corrosion Science 1998;41:463. [73] Tao S, Li DY. Nanotechnology 2006;17:65. [74] Hellmig RJ, Janecek M, Hadzima B, Gendelman OV, Shapiro M, Molodova X, Springer A, Estrin Y. Materials Transactions 2008;49:31. [75] Yu J, Han E, Lu L, Wei X, Leung M. Journal of Materials Science 2005;40:1019. [76] Miyamoto H, Harada K, Mimaki T, Vinogradov A, Hashimoto S. Corrosion Science 2008;50:1215. [77] Li M, Jiang L, Zhang W, Qian Y, Luo S, Shen J. Journal of Solid State Electrochemistry 2007;11:1319. [78] Lu H-B, Li Y, Wang F-H. Thin Solid Films 2006;510:197. [79] Youssef KMS, Koch CC, Fedkiw PS. Corrosion Science 2004;46:51. [80] Wang XY, Li DY. Electrochimica Acta 2002;47:3939. [81] Yu B, Woo P, Erb U. Scripta Materialia 2007;56:353. [82] Luo W, Xu Y, Wang Q, Shi P, Yan M. Corrosion Science 2010;52:3509. [83] Chianpairot A, Lothongkum G, Schuh CA, Boonyongmaneerat Y. Corrosion Science 2011;53:1066. [84] Vinogradov A, Mimaki T, Hashimoto S, Valiev R. Scripta Materialia 1999;41:319. [85] Peng X, Zhang Y, Zhao J, Wang F. Electrochimica Acta 2006;51:4922. [86] Mishra R, Balasubramaniam R. Corrosion Science 2004;46:3019. 151 [87] Ghali E. Corrosion resistance of aluminum and magnesium alloys. New Jersey: Wiley, 2010. [88] Carroll MC, Gouma PI, Mills MJ, Daehn GS, Dunbar BR. Scripta Materialia 2000;42:335. [89] Goswami R, Spanos G, Pao PS, Holtz RL. Materials Science and Engineering A 2010;527:1089. [90] Goswami R, Spanos G, Pao PS, Holtz RL. Metall and Mat Trans A 2011;42A:348. [91] Zhu Y, Cullen DA, Kar S, Free ML, Allard LF. Metall and Mat Trans A 2012;43:4933. [92] Goswami R, Holtz RL. Metall and Mat Trans A 2013;44:1279. [93] Zhu Y. Characterization of beta phase growth and experimental validation of long term thermal exposure sensitization of AA5xxx alloys. Department of Metallurgical Engineering, vol. Master of Science: The University of Utah, 2013. p.116. [94] Gupta RK, Zhang R, Davies CHJ, Birbilis N. Corrosion 2013;69:1081. [95] DoITPoMS. Al-Mg phase diagram. vol. 2014. http://www.doitpoms.ac.uk/miclib/pds.swf?targetFrame=Al-Mg: University of Cambridge. [96] Gao J, Quesnel DJ. Metall and Mat Trans A 2011;42:356. [97] Porter DA, Easterling KE. Phase transformations in metals and alloys. London, UK: Chapman & Hall, 1992. [98] Zener C. J. Appl. Phys. 1949;20:950. [99] Choi D-H, Ahn B-W, Quesnel DJ, Jung S-B. Intermetallics 2013;35:120. [100] Oguocha INA, Adigun OJ, Yannacopoulos S. Journal of Materials Science 2008;43:4208. [101] Yuan Y. Localised corrosion and stress corrosion cracking of aluminium- magnesium alloys. Metallurgy and Materials, vol. Doctor of Philosophy: University of Birmingham, 2006. p.199. 152 [102] Leng Y. Materials characterization: Introduction to microscopic and spectroscopic methods. Singapore: Wiley, 2008. [103] Callister WD. Materials science and engineering: An introduction. York, PA, USA: Wiley, 2007. [104] Brandon D, Kaplan WD. Microstructural characterization of materials. England: Wiley, 2008. [105] Hafner B. Energy dispersive spectroscopy on the SEM: A primer. http://www.charfac.umn.edu/instruments/eds_on_sem_primer.pdf: University of Minnesota - Twin Cities. p.26. [106] miscroanlysis Cffama. Introduction to energy dispersive X-ray spectrometry (EDS). http://micron.ucr.edu/public/manuals/EDS-intro.pdf: University of California Riverside, 2014. p.12. [107] EDAX. Crystallographic orientation. AMETEK Inc., 2002. p.4. [108] Adam J. Schwartz MK, Brent L. Adams, David P. Field. Electron backscatter diffraction in materials science. USA: Springer, 2009. [109] EDAX. Electron backscatter diffraction & orientation imaging microscopy. AMETEK Inc., 2002. p.6. [110] Randle V. Scripta Materialia 2001;44:2789. [111] Wright SI, Larsen RJ. Journal of Microscopy 2002;205:245. [112] Keller RR, Geiss RH. Journal of Microscopy 2012;245:245. [113] Trimby PW. Ultramicroscopy 2012;120:16. [114] Trimby PW, Cao Y, Chen Z, Han S, Hemker KJ, Lian J, Liao X, Rottmann P, Samudrala S, Sun J, Wang JT, Wheeler J, Cairney JM. Acta Materialia 2014;62:69. [115] Volkert CA, Minor AM. MRS Bulletin 2007;32:389. [116] Brydson R. Aberration-corrected analytical transmission electron microscopy. United Kingdom: Wiley, 2011. [117] Farmer M. Electron spectroscopic imaging and electron energy loss spectroscopy (EELS). caur.uga.edu/ppt/EELSWu.ppt: University of Georgia, 2002. 153 [118] GAMRY. Getting started with electrochemical corrosion measurement. Application note: GAMRY Instruments. [119] Robert G. Kelly, John R. Scully, Shoesmith. DW, Buchheit RG. Electrochemical techniques in corrosion science and engineering: CRC Press, 2002. [120] ASTM. Standard practice for conventions applicable to electrochemical measurements in corrosion testing. ASTM standard G3 - 89, Reapproved 2010. [121] Popović M, Romhanji E. Materials Science and Engineering: A 2008;492:460. [122] Jain S, Lim MLC, Hudson JL, Scully JR. Corrosion Science 2012;59:136. [123] Gunson J. Effect of sensitisation on the corrosion fatigue properties of AA5456- H116. Materials and Metallurgy, vol. Masters of Research: University of Birmingham, 2010. p.119. [124] Unocic KA, Mills MJ, Daehn GS. Microscopy and Microanalysis 2008;14:556. [125] Zhao Y, Furnish TA, Kassner ME, Hodge AM. J. Mater. Res. 2012;27:3049. [126] ASTM. Determining average grain size using electron backscatter diffraction (EBSD) in fully recrystallized polycrystalline materials. ASTM Standard E2627. West Conshohocken, PA, 2010. [127] Humphreys FJ. Journal of Materials Science 2001;36:3833. [128] Hodge AM, Wang YM, Barbee Jr TW. Scripta Materialia 2008;59:163. [129] Doherty RD, Hughes DA, Humphreys FJ, Jonas JJ, Juul Jensen D, Kassner ME, King WE, McNelley TR, McQueen HJ, Rollett AD. Materials Today 1998;1:14. [130] Janssens KGF, Olmsted D, Holm EA, Foiles SM, Plimpton SJ, Derlet PM. Nat Mater 2006;5:124. [131] Holm EA, Foiles SM. Science 2010;328:1138. [132] Field DP, Bradford LT, Nowell MM, Lillo TM. Acta Materialia 2007;55:4233. [133] Fullman RL, Fisher JC. J. Appl. Phys. 1951;22:1350. [134] Lingk C, Gross ME, Brown WL. J. Appl. Phys. 2000;87:2232. 154 [135] Berger A, Wilbrandt PJ, Ernst F, Klement U, Haasen P. Progress in Materials Science 1988;32:1. [136] Kobiyama M, Inami T, Okuda S. Scripta Materialia 2001;44:1547. [137] Zhao Y, Cheng IC, Kassner ME, Hodge AM. Acta Materialia 2014;67:181. [138] Fonsati M, Zucchi F, Trabanelli G. Electrochimica Acta 1998;44:311. [139] Vázquez M, Sanchez SRD. J Appl Electrochem 1998;28:1383. [140] Deslouis C, Tribollet B, Mengoli G, Musiani MM. J Appl Electrochem 1988;18:374. [141] Kear G, Barker BD, Walsh FC. Corrosion Science 2004;46:109. [142] Kim SH, Erb U, Aust KT. Scripta Materialia 2001;44:835. [143] Schuh CA, Anderson K, Orme C. Surface Science 2003;544:183. [144] Asgari H, Toroghinejad MR, Golozar MA. Applied Surface Science 2007;253:6769. [145] Chmielová M, Seidlerová J, Weiss Z. Corrosion Science 2003;45:883. [146] Sandberg J, Odnevall Wallinder I, Leygraf C, Le Bozec N. Corrosion Science 2006;48:4316. [147] Núñez L, Reguera E, Corvo F, González E, Vazquez C. Corrosion Science 2005;47:461. [148] Ikemiya N, Kubo T, Hara S. Surface Science 1995;323:81. [149] Kunze J, Maurice V, Klein LH, Strehblow H-H, Marcus P. J. of Phys. Chem. B 2001;105:4263. [150] Kunze J, Maurice V, Klein LH, Strehblow H-H, Marcus P. Corrosion Science 2004;46:245. [151] Du Y, Chang YA, Huang B, Gong W, Jin Z, Xu H, Yuan Z, Liu Y, He Y, Xie F- Y. Mater. Sci. Eng. A-Struct. Mater. Prop. Microstruct. Process. 2003;363:140. [152] Tromans D, Sun R-h. Journal of The Electrochemical Society 1991;138:3235. 155 [153] Rahmouni K, Keddam M, Srhiri A, Takenouti H. Corrosion Science 2005;47:3249.

Abstract (if available)

Abstract Grain boundaries play a significant role in the properties of polycrystalline metals, and the objective of grain boundary engineering is to incorporate special grain boundaries into materials to achieve better properties. While a special grain boundary is conventionally defined as a boundary with coincidence site lattice Σ ≤ 29, it should be noted that the grain boundary plane can also largely affect grain boundary structure and properties. Nanotwins in particular, which are bundles of coherent Σ3 boundaries, have extremely low energy and improved properties compared with ordinary grain boundaries. While it is well‐studied that nanotwins could enhance the mechanical properties of a metal, the effect of nanotwins on the thermal stability and corrosion resistance still needs to be explored and is the subject of this research. In addition to nanotwins which have {111} oriented boundary planes, this study also explored how differently oriented grain boundary planes could affect the sensitization behavior of an Al‐Mg alloy. ❧ For the thermal stability study, sputtered Cu foils with two types of microstructures, highly nanotwinned and non‐nanotwinned columnar grains, were sequentially heat‐treated at 200, 300, and 400℃. The highly nanotwinned Cu (grain size, ∼ 700 nm) remained thermally stable up to 300℃, whereas the non‐nanotwinned Cu (grain size, ∼ 400 nm) had a rapid grain growth even at 200℃. Additionally, the effect of nanotwins on corrosion behavior was evaluated through testing Cu samples containing various fractions of nanotwins: a) highly nanotwinned columnar grains, b) partially nanotwinned columnar grains, c) non‐nanotwinned columnar grains, and d) microcrystalline grains. These samples were tested in 3.5% NaCl solution (pH ~ 8.0) by linear polarization, potentiodynamic polarization, and immersion corrosion methods. It was found that highly nanotwinned Cu had a lower corrosion current density and a more protective passive layer compared with non‐nanotwinned, partially nanotwinned, and microcrystalline Cu samples. The improved thermal stability and corrosion resistance of highly nanotwinned Cu were attributed to the ordered structure of nanotwins, the special grain boundary network, and the {111} texture due to the presence of nanotwins. ❧ While a {111} oriented coherent boundary plane of nanotwins can enhance the thermal and corrosion behavior, this study further identified grain boundary plane orientations and explored their roles in the sensitization behavior of conventional and sputtered Al 5456 samples. Both types of samples were heated at 175℃ and subsequently etched in 10% H₃PO₄ in order to reveal the Mg-rich β phase. The orientations of selected grain boundary planes were measured and related to their corresponding β phase thicknesses. It was found that grain boundaries with plane orientations close to {110} appear particularly vulnerable to β precipitation. Moreover, the sputtered Al 5456 had much higher resistance to β precipitation than the conventional Al 5456, which could be due to its special columnar grain boundary plane orientations. This study highlights the potential of incorporating nanotwins and altering grain boundary plane orientations in order to improve thermal stability, corrosion resistance, and sensitization behavior.

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Asset Metadata

Creator Zhao, Yifu (author)

Core Title The role of nanotwins and grain boundary plane in the thermal, corrosion, and sensitization behavior of nanometals

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Mechanical Engineering

Publication Date 06/24/2014

Defense Date 05/22/2014

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag annealing,corrosion,grain boundary structure,nanotwins,OAI-PMH Harvest,precipitation

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Hodge, Andrea M. (committee chair), Kassner, Michael E. (committee chair), Eliasson, Veronica (committee member), Wang, Pin (committee member)

Creator Email joyful628@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-425150

Unique identifier UC11286599

Identifier etd-ZhaoYifu-2583.pdf (filename),usctheses-c3-425150 (legacy record id)

Legacy Identifier etd-ZhaoYifu-2583.pdf

Dmrecord 425150

Document Type Dissertation

Format application/pdf (imt)

Rights Zhao, Yifu

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA